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Electron irradiation induced amorphous SiO2 formation at metal oxide/Si interface at room temperature; Electron beam writing on interfaces

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Al2O3 (5 nm)/Si (bulk) sample was subjected to irradiation of 5 keV electrons at room temperature, in a vacuum chamber (pressure 1 × 10-9 mbar) and formation of amorphous SiO2 around the interface was observed. The oxygen for the silicon dioxide growth was provided by the electron bombardment induced bond breaking in Al2O3 and the subsequent production of neutral and/or charged oxygen. The amorphous SiO2 rich layer has grown into the Al2O3 layer showing that oxygen as well as silicon transport occurred during irradiation at room temperature. We propose that both transports are mediated by local electric field and charged and/or uncharged defects created by the electron irradiation. The direct modification of metal oxide/silicon interface by electron-beam irradiation is a promising method of accomplishing direct write electron-beam lithography at buried interfaces.
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Electron irradiation induced
amorphous SiO2 formation at
metal oxide/Si interface at room
temperature; electron beam
writing on interfaces
S. Gurbán1, P. Petrik1, M. Serényi1, A Sulyok1, M. Menyhárd
1, E. Baradács2, B. Parditka2,
C. Cserháti2, G. A. Langer2 & Z. Erdélyi2
Al2O3 (5 nm)/Si (bulk) sample was subjected to irradiation of 5 keV electrons at room temperature, in
a vacuum chamber (pressure 1 × 1 09 mbar) and formation of amorphous SiO2 around the interface
was observed. The oxygen for the silicon dioxide growth was provided by the electron bombardment
induced bond breaking in Al2O3 and the subsequent production of neutral and/or charged oxygen.
The amorphous SiO2 rich layer has grown into the Al2O3 layer showing that oxygen as well as silicon
transport occurred during irradiation at room temperature. We propose that both transports are
mediated by local electric eld and charged and/or uncharged defects created by the electron
irradiation. The direct modication of metal oxide/silicon interface by electron-beam irradiation is a
promising method of accomplishing direct write electron-beam lithography at buried interfaces.
Due to the paramount role of SiO2 in integrated circuits its formation is extremely well studied and understood.
e really good quality silicon dioxide is produced at elevated temperatures to cope with the activation energy of
its formation. It is well known, however, that even at room temperature native oxide forms at the surface of Si. e
growth of native oxide terminates at a thickness of 2–3 nm. It has also been shown that SiO2 growth at the surface
of Si at low temperatures greatly facilitated if the oxygen molecule is excited by some means. is excitation can
be carried out by various ways e.g. various plasma processes1, noble-gas ion bombardment2 electron irradiation3,4,
etc. Much less is known about the formation of SiO2 at metal oxide/silicon interface.
Aluminum oxide/Si interface has been frequently studied for at least two reasons: a./in a quest of high die-
lectric constant material to replace the SiO25, b./for passivation of the Al2O3/Si interface in photovoltaic applica-
tions69 and quantum dots10. In both cases, the presence of xed charges inherent to Al2O3 causes a problem11,
which can be addressed by various means12. e interface of the as deposited (chemical vapor deposition at
400 °C) system of a 3.5 nm thick Al2O3 lm on Si(100) was studied by Klein et al.13, using nuclear reaction res-
onance proling. ey could detect the presence of a silicate layer. e theoretical calculation of Xiang et al.14
showed the existence of a sharp Al2O3/Si interface exhibiting Si-O and Si-Al covalent bonds. us, in these works
Si-O bond formation has been observed at the interface, but SiO2 formation has not been reported.
On the other hand medium energy electron and/or X ray irradiation might cause various chemical changes
(including oxidation), defect formation, structural changes in the surface close regions of the irradiated material,
mainly in the case of insulators, where recombination and healing are limited. Defect formation15,16 is explained
by bond breaking, while structural changes like amorphization15,17 is explained by excitation of the chemical
bonds and subsequent atomic movements.
In this paper, we will report on the electron irradiation induced amorphousSiO2 growth at the Al2O3/Si inter-
face. To demonstrate this process samples were made by growing a 5 nm thick Al2O3 layer onto a carefully cleaned
1Institute for Technical Physics and Materials Science, Centre for Energy Research Hungarian Academy of Sciences,
P.O.B. 49, H-1525, Budapest, Hungary. 2Department of Solid State Physics, University of Debrecen, P.O. Box 400,
H-4002, Debrecen, Hungary. P. Petrik, M. Menyhárd and Z. Erdélyi contributed equally to this work. Correspondence
and requests for materials should be addressed to M.M. (email: menyhard.miklos@energia.mta.hu)
Received: 13 November 2017
Accepted: 19 January 2018
Published: xx xx xxxx
OPEN
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(100) Si substrate by two methods, by atomic layer deposition (ALD) and by radio frequency (RF) sputtering. e
samples were then irradiated by electrons in a vacuum system (pressure of 109 mbar) at room temperature. e
parameters of electron irradiation were: energy 5 keV, current density up to 6 × 1016 electrons/cm2/s, total uence
up to 6 × 1021 electrons/cm2. e vacuum system was equipped with electron and ion guns and an electron ana-
lyzer. e eect of irradiation was studied in situ by Auger Electron Spectroscopy (AES), AES depth proling, and
ex situ by spectroscopic ellipsometry (SE), and transmission electron microscopy (TEM).
Results
Transmission Electron Microscopy (TEM). TEM images were taken before and aer irradiation; they did
not show any crystalline phase except the Si substrate. us, both phases, Al2O3 and SiO2, observed were amor-
phous. To emphasize that the SiO2 produced is amorphous, we will sign it as a: SiO2.
Spectroscopic ellipsometry (SE). e irradiated spot can clearly be identied on the map of measured
ellipsometric angles, Δ, in Fig.1, revealing that the irradiation caused a signicant change in the optical proper-
ties of the layers. e dierence in Δ between the irradiated and non-irradiated regions is larger than one degree,
thus more than an order of magnitude higher than the sensitivity of the measurement.
For determination of the thickness of the amorphous SiO2 (a:SiO2) the interface was modeled as an eective
medium composition of Al2O3, SiO2 and Si18. Figure2 shows that both the interface thickness, a, and the volume
fraction of a:SiO2, b, increase in the irradiated spot. Due to small thickness and the correlation between the
Al2O3 and the a:SiO2 components, there might be a signicant error in the determined volume fraction of a:SiO2.
However, the trend of the increasing interface thickness and volume fraction of a:SiO2 in the irradiated part is
very clear from the ellipsometry results.
Auger Electron Spectroscopy (AES). The samples were produced either by ALD or RF sputtering.
e results did not depend on the production technology and thus in the following the type of sample will
not be specied. Figure3 shows the Auger spectra, N(E), in the 1600 ± 20 eV range (where the SiKLL Auger
line appears independently from its chemical state) recorded on the surface of the irradiated (1.32 electrons/s/
A2×20 hours = 9.4×1020 electrons/cm2) and non-irradiated regions of the sample. It is clear that the intensity
of the KLL Auger electrons emitted by Si in oxide environment increased considerably due to the irradiation.
Figure 1. Map of the measured ellipsometric angles, Δ, around the illuminated spot at the wavelength of
300 nm.
Figure 2. Maps of thickness of the interface layer, (a) and the volume fraction of a:SiO2 in the interface layer, (b).
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Simultaneously, the intensity of the KLL Auger electrons emitted by Si in metal environment decreased. e
rough interpretation of the nding is simple; within the information depth of the SiKLL Auger electrons a fraction
of the Si has been converted to SiOx. e x value cannot be accurately determined by AES but it is not far from 2.
From the TEM results it is also clear that the compound formed is amorphous. us we will sign the compound
as a:SiO2, which means an amorphous phase with composition close to the stoichiometric. However, since the
measured Auger intensity is a weighted sum of the intensities emitted from various depths, the actual depth dis-
tribution of the a:SiO2 produced cannot be derived from this measurement.
For the determination of the depth distribution of the a:SiO2 produced by the electron irradiation AES depth
proling has been utilized. In AES depth proling the surface of the sample is etched away in predened steps
by using ion sputtering and the newly exposed surface is analyzed by AES which provides elemental and some
chemical information.
Figure4 shows the raw data (measured Auger peak-to-peak intensities in the dierentiated, N(E) curve)
obtained from the non-irradiated (a) and irradiated (b) (total electron dose of 9.4×1020 electrons/cm2) regions
of the sample; the two depth proles show considerable dierences. e most important one is that while in
the case of the non-irradiated region the intensity of the Si Auger signal in oxide environment (signed by SiO)
is just higher than the noise level around the Al2O3/Si interface, in the case of the irradiated region the same
signal is easily measurable, proving that due to the irradiation a:SiO2 has formed. It should also be noted that
longer sputtering time is necessary to remove the layer at the irradiated region than that necessary in the case
of the non-irradiated one, demonstrating that the thickness of the top most layer in the irradiated region (pure
Al2O3 + altered layer + a:SiO2) of the sample is greater than that of the virgin part. Figure5 shows the concentra-
tion distributions, which best correspond to the depth proles shown in Fig.4 measured on the non-irradiated
(virgin) and irradiated region of the sample, respectively, calculated by our trial and error evaluation method19,20.
Note that the irradiation took place at room temperature.
Figure 3. e raw data, measured intensity (N(E)) vs energy (E) in the vicinity of the Si KLL transition obtained
from the non-irradiated region of the surface (non-irrad.) and aer an irradiation of 9.4×1020 electrons/cm2
(irrad.).
Figure 4. e raw data, peak-to-peak intensities measured in the N(E) curve, of AES depth proling obtained
from non-irradiated (a) and irradiated (b) regions of the sample. SiO and AlO mean Auger intensities of Si and
Al being in oxygen environment, respectively.
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Figure5a shows that the structure of the non-irradiated sample is the expected one; there is an about 5 nm
thick Al2O3 layer on the top of the Si substrate. e small amount of a:SiO2 which is at the interface region is due
to either the non-proper cleaning of the Si substrate or oxidation at the beginning of the deposition of the Al2O3.
On the other hand, the irradiated region provides a strongly altered structure. ough the surface is covered by a
pure Al2O3 layer its thickness is only about 3 nm instead of the 5 nm initial thickness. is cover layer is followed
by an Al2O3 and a:SiO2 mixture in which the concentration of the a:SiO2 increases toward the interface. e over-
all thickness of the altered layer and reminder Al2O3 layer is larger than that of the initial Al2O3 layer.
Discussion
Both methods (SE and AES) clearly show that the irradiated part of the sample has considerably changed: a:SiO2
has been built due to the irradiation. Since the non-destructive ellipsometric measurement cannot produce a:SiO2
it is clear that the a:SiO2 is the result of the electron irradiation.
An obvious question is whether the subsequent AES studies combined with ion bombardment might impair
the actual distribution of the a:SiO2 formed during the electron beam irradiation experiment. Considering Fig.5b
we can estimate the overall rate of the process; it is rather low being in the range of 104–105 a:SiO2/5 keV elec-
trons. us, the production of SiO2 during AES measurement can be ignored since the irradiation uence during
the recording of the AES spectra is 300 times lower than that of the irradiation experiment.
e ion bombardment can also produce SiO2. e eciency of the process can also be easily checked exper-
imentally. e measured Auger intensities recorded on the surface before Auger depth proling cannot be used
to determine the distribution of the emitting elements. On the other hand, the depth distributions determined
by the AES depth proling can easily be integrated and they should give the intensities measured at the surface
before the AES depth proling. is integral has been calculated for all measurements and compared with the
intensities measured on the surface. e amount of a:SiO2 calculated from the AES measurement performed on
the surface before the AES depth proling started, and that obtained for the AES depth prole agreed within an
error of 20%, meaning that the possible eect of the ion bombardment on the a:SiO2 content is typically stillless
than 20%, which will be ignored. us, the concentration distributions provided by the AES depth proling can
be accepted as those produced by the 5-kV electron irradiations.
e accurate determination of heating eect of the electron beam irradiation is rather dicult15 in case of high
current densities and small particles. In our case the current density was chosen to be rather low, which allows to
make a rough overestimation of the temperature rise, ΔT. It can be supposed that the whole irradiation energy
is absorbed by the sample and used for heating. In this case in stationary case the input energy is equal to that
carried away by heat conduction which results in the simple equation of ΔT = Pr/K21, where P is the irradiation
power density, r is the radius of the electron beam, and K is the heat conduction. In the present case ΔT is around
3 K, which is negligible, and consequently all processes are taking places at room temperature.
e basic features of the a:SiO2 growth (due to the electron irradiation) were similar on samples prepared by
dierent technologies. us, we ignore possible eects of the imperfections connected to the sample production
technology, and the process will be explained assuming perfect interfaces. e growth of a:SiO2 can be divided
into two steps: creating charged and/or neutral oxygen and the growth process itself.
Electron irradiation induced bond breaking & excitation. Vast amount of experimental data has been
amassed in transmission electron microscopy (TEM) studies concerning various electron beam-induced dam-
ages in insulators15. Several electron irradiation-induced eects, including phase transformation, decomposition,
amorphization, oxidation, reduction, etc., have been identied. For the description of the phenomena two basic
models have been developed: a./the mechanical interaction of the bombarding electron with the nuclei, knock-on
mechanism, b./the electronic excitation by the electric eld, radiolytic processes.
e knock-on process in our case will not be considered since the energy, 5 keV, of our bombarding electrons
is far below the threshold of this mechanism15. On the other hand, the radiolytic processes have practically no
Figure 5. e concentration distributions of the sample at the non-irradiated (a) and irradiated (b) regions,
respectively derived from the depth proles shown in Fig.4a and b.
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threshold since because of the multiple interactions nearly any type of excitations can happen. E.g., recently
the dissolution of boehmite under the high energy electron irradiation (TEM studies) was explained by the
electron-hole pairs created during the electron bombardment22. is explanation is rather similar to those devel-
oped long time ago for the description of low-energy electron-stimulated desorption (ESD)2325 which can readily
be applied for the explanation of our observations. According to these models following the ionization of a core
electron, a valence electron from the O decays to ll the resulting hole in frame of the Auger process. erefore,
O2 transforms to O0. It might also happen that by an additional Auger eect a further electron is emitted and
O+ is formed for a short while2325. e lifetime of both O0 and O+ is long, since fast recombination is not pos-
sible due to the insulator matrix. During the excitation processes, because of Columbic repulsion, structural
changes might also occur. e electron irradiation this way produces neutral and/or charged O in the Al2O3
matrix together with various charged and neutral crystal defects.
SiO2 growth. If Si surface containing O2 and/or H2O is irradiated by electrons, SiO2 can be produced3,4. e
description of the process is similar to that of ESD. In this case it is assumed that electrons attach to an adsorbed
O2 molecular precursor to form O2. e O2 then decomposes to form O and O, and one or both of these spe-
cies cause rapid oxidation of the surface. In our case neutral and/or charged O are produced in the Al2O3 matrix
and for the compound formation to take place they should be transported to the interface.
e diusion of O in α alumina has been studied by Sokol et al.26 showing, that due to the structural and spin
conguration, the defect reaction energy can change by over 2 eV. is behavior aects the equilibrium defect
concentrations by many orders of magnitude. Consequently, the diusion processes in such materials may be
more complicated, which has previously been assumed. Århammar et al. reached similar conclusion for amor-
phous oxides27. Still it seems that at room temperature neither the O nor the Si can be transported fast enough by
thermal diusion to explain the growth of SiO2 layer at the Al2O3/Si interface.
It is well known that despite the high heat of the SiO2 formation the native oxide forms on the surface of clean
Si at room temperature in air. e thickness of this oxide is about 2 nm. is process is explained by the theory
of Cabrera and Mott, which assumes that the electrons can pass freely from the Si to the oxide surface to ionize
oxygen atoms. is establishes a uniform eld within the oxide, which leads to a shi in the Fermi level of the
oxide28. e same reasoning can be used if electron is placed to the surface from any other source. Nowak et al.
have shown that electron bombardment, providing charges to build up electric eld, induces oxide growth on
tungsten nanowires at room temperature29,30. ey use the explanation of Mott and Cabrera; the electric eld
created reduces the energy barriers for the migration of metal cations or oxygen anions into and through the
oxide, allowing signicant material transport and thus growth of the oxide layer at low temperature. It is also evi-
dent that this is a self-controlled process; aer reaching a given thickness the strength of electric eld is no more
sucient to drive the diusion.
Our Al2O3 layer is 5 nm thick that is thicker than the native SiO2. In our case, however, the charged defects,
charged and neutral oxygen produced by the electron bombardment are distributed evenly in the layer. at is, in
our case various distributions of local elds and charged as well as neutral oxygen is produced. It is evident that
those oxygen atoms, which are close to the interface can be transported to the substrate Si atoms and form oxide
at a high probability; thus the oxide growth starts from the interface. Similarly, if the charged defect is close to the
interface, then an electric eld of sucient strength is built up initiating the diusion of Si into the Al2O3 laye r.
Since in the defected Al2O3 layer there are charged and/or neutral oxygen atoms, compound formation might
take place. e a:SiO2 grain grows by adding additional Si and O atoms. e primary Si source is the substrate.
It should be remembered, however, that the self-limitation process is activated aer the a:SiO2 grain reaches a
certain thickness and the local eld strength is not sucient anymore to drive the diusion. It seems that our
particles are somewhat larger than that of the typical thickness of the native oxide. On the other hand, one should
also consider that the already produced a:SiO2 grains are also subjected to electron irradiation producing charged
defects, excited oxygen atoms and quasi free Si in the a:SiO2 grain. e quasi free Si can utilize the local elds to
diuse to the surface of the a:SiO2 grain resulting in further growth. e probability of this process is lower than
that of the primary one, however, and the rate of growth in this phase is much lower resulting in only some addi-
tional growth resulting in a 3–4 nm thick a:SiO2 grains.
Conclusions
We have shown that bombarding 5 nm thick Al2O3/Si structure by 5 keV electrons at room temperature, amor-
phous silicon dioxide is produced. The amorphous SiO2 grains grow from the interface toward the Al2O3
matrix; their amount depends on the irradiated charge. e phenomenon was explained considering electron
bombardment-induced bond breaking in Al2O3, electric eld driven diusion of Si and O in defected regions of
the Al2O3 based on the Cabrera-Mott theory.
e direct modication of metal oxide/silicon interface by electron irradiation is a promising method of
accomplishing direct write lithography at buried interfaces.
Samples and Methods. Samples. Samples were made by growing an Al2O3 layer on a Si (100) substrate
ALD and RF sputtering.
e ALD layers were made by a Beneq TFS 200 ALD reactor in the plasma-enhanced deposition mode.
Trimethylaluminium (TMA – from Sigma Aldrich) together with high purity oxygen gas was used for deposition.
Prior to the sample preparation, the deposition chamber had been heated up to 150 °C. During the deposition, the
pressure inside the vacuum chamber was 9.5 mbar, while in the reactor chamber 1.1 mbar. e RF power of the
plasma was set to be 50 W and the ow rate of the oxygen was 100 sccm. e ALD cycle was the following: 150 ms
TMA, which was followed by a 2s purge, then a 2s oxygen plasma at 50 W and at the end of the cycle another 2s
purge.
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e RF sputter deposition was carried out in a Leybold Z400 apparatus evacuated to 5×105 mbar. Sputtering
was performed under a mixture of high purity argon and oxygen gases with an applied RF power of 255 W yield-
ing a plasma pressure of 2.5×102 mbar. Oxygen was continuously let into the sputtering chamber at ow rates
resulting in a partial oxygen pressure of 6%. e deposited amorphous Al2O3 lm has a high refractive index and
low absorption coecient31.
Electron irradiation and Auger Electron Spectroscopy. All electron irradiation experiments have been carried out in
our standard vacuum system equipped with electron energy analyzer (STAIB OPC 105) and various electron sources.
e electron irradiation was facilitated by an electron gun (STAIB EK-10-M) with the following parameters: energy
5 keV, current density up to 6*1016 electrons/cm2/s uence about up to 6 × 1021 electrons/cm2, angle of incidence 54°
(with respect to surface normal). e irradiated area was 100 × 100 μm2. For AES analysis the same electron gun was
used for the excitation with a current density and energy of 1*1016 electrons/cm2/s and 5 keV, respectively. e uence
during the recording of the Auger spectra is about 300 times less than that used for irradiation.
e Auger spectra, N(E), were recorded in counting mode. e recorded spectrum was numerically dieren-
tiated, N(E), for performing the concentration calculation.
e Auger signals of AlKLL, AlLVV, SiKLL, SiLV V, C and O, were measured; by measuring the high (KLL) and low
(LVV) escape depths Auger electrons the quality of evaluation of the depth distributions in AES depth proling
considerably improves. e energies of the Al and Si Auger electrons strongly depend on their chemical environ-
ment; the corresponding Auger electron energies (in eV) are shown in Table1.
erefore it is easy to determine the metal and oxide fractions of the elements by measuring either the LVV or
KLL Auger electrons allowing the determination of the depth distributions for the metal and oxide components
separately.
AES depth proling. A low energy ion gun of Technoorg Linda was used for AES depth proling. e param-
eters of the ion bombardment used were: energy 1 keV, projectile Ar, angle of incidence (with respect to the
surface normal) 80° and specimen rotation during ion bombardment. e ion beam was scanned in an area of
1.5 × 1.5 mm2. Using these parameters, the ion bombardment induced roughening and mixing is minimal32.
Determination of the concentration distribution from AES spectra. In AES, generally simple expressions are
applied for the evaluation of the composition using the measured peak-to-peak amplitudes of the dierentiated,
N(E), curve33. is expression assumes that the composition within the escape depth of signal electrons is homo-
geneous; in any other case it cannot be applied. is is however the situation presently since the thickness of the
Al oxide layer is only 5 nm, while the inelastic mean free path (IMFP) of the AlKLL and SiKLL Auger electrons in
Al2O3 are 3.2 and 3.6 nm, respectively34. We used a trial and error approach to determine the composition distri-
bution of our sample19,20. e essence of the method is that we assume a composition distribution along the depth
and calculate the Auger intensities assuming that the transport of electrons can be described by the exponential
attenuation law, not considering the elastic scattering. (Neglecting the elastic scattering creates an error in the
range of 10–15%, which will not aect the description of the phenomena.) e composition distributions are var-
ied until the simulated depth prole is close enough to the measured one. If an element emits high (high IMFP)
and low energy (low IMFP) Auger electrons, as in the present case, the accuracy of the method is rather good.
In case of depth proling the above procedure is repeated for all spectra obtained aer each sputtering steps
assuming that the ion bombardment used for removing the material does not cause serious changes to the mate-
rial. is is a reasonable assumption since the removed layer thickness is less than 8 nm and all alterations scale
with the removed layer thickness32.
Spectroscopic ellipsometry (SE). Auger depth proling uses ions and electrons to reveal the depth distribution of
the SiO2. Both projectiles may initiate the formation of SiO2. ough it will be shown that these are low probabil-
ity processes, still we have applied SE, a non-destructive method, to verify the presence of the SiO2 produced by
electron irradiation. e SE measurements have been carried out by a Woollam M-2000DI rotating compensator
ellipsometer at an angle of incidence of 55°. e microspot option was used to focus the light into a spot with
a diameter of approximately 0.2 mm. e surface around the irradiated region was mapped with steps smaller
than the spot size, in order to precisely locate the modied spot on the sample surface, and to measure only the
irradiated region of the sample.
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LVV KLL
Oxide Metal Oxide Metal
Al 54 68 1389 1396
Si 78 92 1610 1619
Table 1. e LVV and KLL Auger lines energies (eV) of Al and Si in pure and oxidized forms.
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Acknowledgements
e work is supported by the GINOP-2.3.2-15-2016-00041, NN 114422 M-ERANET “GRACE” and TÉT_16-1-
2016-0100 projects. e project is co-nanced by the European Union and the European Regional Development
Fund.
Author Contributions
Z.E. and M.M., co-wrote the paper, headed and coordinated the work. E.B. and B.P. produced the ALD layers.
C.Cs. and G.A.L. performed preliminary e-beam irradiation and quality test measurements. M. S. made the layers
by RF sputtering. P.P made the ellipsometry measurements and evaluation and co-wrote the paper. E.B., C. Cs.,
G.A.L., A. S. contributed to the discussion of the results. S.G., A.S., M.M. performed the AES studies.
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