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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:
Received: 13 November 2017
Accepted: 19 January 2018
Published: xx xx xxxx
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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).
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
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.
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.
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
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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
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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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
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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... In our previous paper 8 we have studied effect of electron bombardment on the Al 2 O 3 /Si substrate system at room temperature; in this paper we study the same system at various elevated temperature. A great variety of alterations occur when varying the sample temperature during irradiation, ranging from serious degradation of the sample (at 500 °C) to slight metallic Si diffusion to the nearly perfect Al 2 O 3 layer (at 700 °C). ...
... Irradiation at room temperature. The effect of irradiation of the 5 nm Al 2 O 3 /Si substrate sample have been described in detail 8 . Since the effects of irradiation at elevated temperature were expected to be much stronger than that at room temperature, thus we also used thicker, 20 nm Al 2 O 3 /Si substrate sample. ...
... Still thermally activated Si diffusion occurs. Here we note that Si movement was also observed at room temperature since SiO 2 formation occurred, that movement was, however, not thermally activated diffusion but following the Cabrera-Mott process (see also ref. 8). To explain the observed diffusion, we recall that defect production is temperature independent thus defect formation occurs during irradiation at 700 °C. ...
Full-text available
Interface induced diffusion had been identified in a thin film system damaged by electron bombardment. This new phenomenon was observed in Al 2 O 3 (some nm thick)/Si substrate system, which was subjected to low energy (5 keV) electron bombardment producing defects in the Al 2 O 3 layer. The defects produced partially relaxed. The rate of relaxation is, however, was different in the vicinity of the interface and in the "bulk" parts of the Al 2 O 3 layer. This difference creates an oxygen concentration gradient and consequently oxygen diffusion, resulting in an altered layer which grows from the Al 2 O 3 /Si substrate interface. The relative rate of the diffusion and relaxation is strongly temperature dependent, resulting in various altered layer compositions, SiO 2 (at room temperature), Al 2 O 3 + AlO x + Si (at 500 °C), Al 2 O 3 + Si (at 700 °C), as the temperature during irradiation varies. Utilizing this finding it is possible to produce area selective interface patterning.
... In our previous paper 5 we have studied effect of electron bombardment on the Al2O3 / Si substrate system since this structure is frequently applied in photovoltaic applications 6,7 . It was observed that due to room temperature electron bombardment, a SiO2 layer had grown on the Al2O3 / Si interface. ...
... The results of irradiation of 5 nm Al2O3 / Si substrate sample have been reported 5 . Since the effects of the irradiation at elevated temperature are much stronger than that at room temperature, besides the previously applied 5 nm Al2O3 / Si substrate sample, we also used thicker, 20 nm Al2O3 / Si substrate sample. ...
... Still thermally activated Si diffusion occurs. Here we note that Si movement was also observed at room temperature since SiO2 formation occurred, that movement was, however, not thermally activated diffusion but by the Cabrera-Mott process; see also ref. 5. To explain the observed diffusion, we recall that defect production is temperature independent thus defect formation occurs during irradiation at 700 o C. According to eq. (1) to reach the stationary state the presence of some defects is necessary, thus the Al2O3 layer should contain defects. ...
Full-text available
Interface induced diffusion had been identified in thin film system damaged by electron bombardment. This new phenomenon was observed in Al 2 O 3 (some nm thick) / Si substrate system, which was subjected to low energy (5 keV) electron bombardment producing defects in the Al 2 O 3 layer. The defects produced partially relaxed. The rate of relaxation is, however, different in the surrounding of the interface and in the "bulk" parts of the Al 2 O 3 layer. This difference generates an oxygen concentration gradient and consequently oxygen diffusion, resulting in an altered layer which grows from the Al 2 O 3 / Si substrate interface. The relative rate of the diffusion and relaxation is strongly temperature dependent, resulting in various altered layer compositions, SiO 2 (at room temperature), Al 2 O 3 + AlO x +Si (at 500 o C), Si(at 700 o C), as the temperature during irradiation varies. Utilizing this finding it is possible to make area selective interface patterning.
... It can be thanks to the loss of core Fe3O4, which contributes considerably to the absorption [66]. The expansion of SiO2 on the water's surface at low temperatures is greatly expedited if the O2 molecules area unit excited [68]. ...
The existence of water on earth is very abundant and has a vital role in the source of life for every living creature. In managing water resources, pollution is one of the issues world researchers face. This article reviews the characteristics and methods of synthesizing Fe2O3 and SiO2 materials to prevent water pollution. The strategies administrated antecedently square measure vapor deposition, microemulsion, solvothermal, coprecipitation, sol-gel, and hydrothermal. The formation of fine quality nanoparticles with controlled size associate degreed size distribution square measure typically achieved by selecting an applicable solvent mixture and varied parameters like temperature, pressure, and time interval.
... In case B, in HV, in UHC condition, when the SEM and the samples are both clean, no contamination appears. Note that the center of the Si sample gets slightly brighter due formation of Si oxide [5] in the presence of water on the surface. Low-energy electron irradiation-based cleaning of the SEM sample chamber and/or its sample is an effective technique with advantages over other cleaning techniques used for the elimination of precursor molecules of electron-beam-induced contamination. ...
... Moreover, two primary models proposed for justifying the amorphization process under the TEM e-beam irradiation are (a) the mechanical thrust of the bombarding electrons delivered to the target atomic nuclei, i.e., knock-on mechanism; (b) electronic excitation by the electric field, i.e., radiolytic process[269].The physical understanding of the amorphization and its corresponding reasoning, with a list of few specific materials investigated earlier includes: (a) electronic excitation intitanate pyrochlores (anion disorder in A 2 Ti 2 O 7 , A=Y, Gd, and Sm), colloidal CsPbBr 3 nanocrystals, complex rare earth (Y(III), Eu(III), and La(III)) nanostructures etc., whereas (b) the frequently encountered knock-on displacements mechanism are observed in-SiC respectively [270]-[273]. More recently, fast and reversible phase transitions (amorphous to crystalline) have been observed in the case of the chalcogenide phase-change material (i.e., on GeSb 2 Te 4 thin films) examined under TEM e-beam irradiation [274]. ...
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The main objectives of the thesis are 1. Investigate biomineralization mimicked nc-ceria individual building units utilized to grow 1D-ceria nanostructures. In this context, de-ionized water (DI-water) having transparent nc-ceria colloidal dispersion with aging leads to anisotropic 1D-ceria nanorods development. Thereby DI-water as an enabler of NCG, similar to aquatic biominerals, is established. To physically demonstrate the building units’ attachment, in-situ TEM e-beam is utilized to act as a probe and also as an enabler of the NCG scheme. 2. Apart from NCG, a significant section is also devoted to investigating the autocatalytic regenerative surface feature of nc-ceria in delivering a tunable charge transfer (CT) visible photoluminescence (PL) emission and quenching. 3. The feasibility of a variety of ultrasonic-sonochemical physio-chemical processes for nanoscience is tracked. Their mechanistic features are accessed, and also utilization in solution-phase material processing is realized. 4. To explore and present ultrasonic-sonochemical process intensification while evaluating the nucleation, growth, and stabilization aspect of the ultrafine oxide-free PVP embedded Al-rich, Al/PVP composite fuel. 5. Lastly, to investigate physically mixed stoichiometric 1:1 NEM made out of the Al/PVP composite fuel and nc-ceria oxidizer. The NEM characteristics evaluated are oxidation, ignition, the energy contained, and also energetics, respectively. The thesis is organized into six chapters, and the contents of each chapter are described briefly in the following sections.
The Mo5O14-type structure is representative of the MoO-based catalyst in the selective oxidation process. Single-crystalline Mo5O14 nanowires can be synthesized in a controlled manner by chemical vapor deposition (CVD). A nanowire catalyst with a porous structure combines the advantages of both nanoparticles and nanowires, leading to a substantial increase in the specific surface area. Therefore, we aim to manipulate the e-beam irradiation process on Mo5O14 nanowires to induce the nanoporous structures in selected regions. In situ transmission electron microscopy (TEM) enabled us to visualize the structural transformation through gradual e-beam irradiation. The e-beam irradiation process removes oxygen atoms and renders the internal structure unstable. After the irradiated region is exposed to air, atoms tend to escape to decrease the internal energy. This results in the formation of nanopores because of the irradiation effect. By nanoscale modification method, the irradiated region is controlled by the electron beam size, which determines the nanopore distribution in the selected region. The study is beneficial for increasing the surface area of Mo5O14-type catalysts with variable nanopore densities and for modifying nanomaterials using a convenient method.
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A hydrothermally synthesized monoclinic phase of BiEuWO 6 photocatalyst nanoparticles was irradiated with variable doses of high-energy electron beam irradiation (EBI). Structural and morphological changes in unirradiated (referred to as control) and electron-beam-irradiated (referred to as EBI) nanoparticles were characterized by X-ray diffraction, Raman spectroscopy, and field emission scanning electron microscopy. At 50 kGy dose, evolution of a small fraction of a crystalline secondary orthorhombic phase along with a primary monoclinic phase is observed. This is further confirmed by X-ray diffraction as well as Raman spectroscopy. Single-phase and multiphase Rietveld refinements were carried out on the powder X-ray data of the control and EBI samples, and the phase fractions were deduced. Further diffused reflectance spectroscopy, steady-state fluorescence emission spectroscopy, and Brunauer-Emmett-Teller surface area were used to characterize the samples. A significant increase in the visible light photocatalytic activity is observed in the two-phase nanomaterials above an optimum dose of 50 kGy for the degradation of Congo red dye. The structural and morphological implications are investigated in detail to understand the enhancement in the photocatalytic activity of the EBI samples. This work demonstrates the potential of high-energy electron beam irradiation for development of superior crystalline semiconductor photocatalysts.
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Layered (oxy) hydroxide minerals often possess out-of-plane hydrogen atoms that form hydrogen bonding networks which stabilize the layered structure. However, less is known about how the ordering of these bonds affects the structural stability and solubility of these minerals. Here, we report a new strategy that uses the focused electron beam to probe the effect of differences in hydrogen bonding networks on mineral solubility. In this regard, the dissolution behavior of boehmite (γ-AlOOH) and gibbsite (γ-Al(OH)3) were compared and contrasted in real time via liquid cell electron microscopy. Under identical such conditions, 2D-nanosheets of boehmite (γ-AlOOH) exfoliated from the bulk and then rapidly dissolved, whereas gibbsite was stable. Further, substitution of only 1% Fe(III) for Al(III) in the structure of boehmite inhibited delamination and dissolution. Factors such as pH, radiolytic species, and knock on damage were systematically studied and eliminated as proximal causes for boehmite dissolution. Instead, the creation of electron/hole pairs was considered to be the mechanism that drove dissolution. The widely disparate behaviors of boehmite, gibbsite, and Fe-doped boehmite are discussed in the context of differences in the OH bond strengths, hydrogen bonding networks, and the presence or absence of electron/hole recombination centers.
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We report on the structural evolution of tunneling oxide passivating contact (TOPCon) for high efficient solar cells upon thermal annealing. The evolution of doped hydrogenated amorphous silicon (a-Si:H) into polycrystalline-silicon (poly-Si) by thermal annealing was accompanied with significant structural changes. Annealing at 600 °C for one minute introduced an increase in the implied open circuit voltage (Voc) due to the hydrogen motion, but the implied Voc decreased again at 600 °C for five minutes. At annealing temperature above 800 °C, a-Si:H crystallized and formed poly-Si and thickness of tunneling oxide slightly decreased. The thickness of the interface tunneling oxide gradually decreased and the pinholes are formed through the tunneling oxide at a higher annealing temperature up to 1000 °C, which introduced the deteriorated carrier selectivity of the TOPCon structure. Our results indicate a correlation between the structural evolution of the TOPCon passivating contact and its passivation property at different stages of structural transition from the a-Si:H to the poly-Si as well as changes in the thickness profile of the tunneling oxide upon thermal annealing. Our result suggests that there is an optimum thickness of the tunneling oxide for passivating electron contact, in a range between 1.2 to 1.5 nm.
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In this Report we show the role of charge defects in the context of the formation of electrostatically defined quantum dots. We introduce a barrier array structure to probe defects at multiple locations in a single device. We measure samples both before and after an annealing process which uses an Al2O3 overlayer, grown by atomic layer deposition. After passivation of the majority of charge defects with annealing we can electrostatically define hole quantum dots up to 180 nm in length. Our ambipolar structures reveal amphoteric charge defects that remain after annealing with charging energies of 10 meV in both the positive and negative charge state.
Amorphous aluminum oxide (a-Al2O3, alumina) can be widely used for ceramic coatings, gate oxide for microelectronics and waveguiding component of integrated optical elements. Moreover, it is a candidate for masks and molds for the preparation of new generation nanoscale devices. Among different technological procedures, cathode sputtering is one of the most effective techniques to deposit amorphous materials which could not be vitrified by an ordinary melting method. Here, the structural and optical properties of Direct Current (DC) magnetron and Radio Frequency (RF) sputtered alumina layers have been revealed regarding to the preparation method. It is shown that the optical absorption and the refractive index of the RF sputtered alumina enable the films to be used as high quality waveguiding material. The oxygen incorporation from the plasma with higher oxygen content results in a bandgap shift to the lower values. Contrarily, reactive DC magnetron sputtering process led to only partly oxidized film growth exhibiting higher absorption.
Plasma-enhanced chemical vapor-deposited (PECVD) silicon oxide films were prepared by the reaction of SiH//4 and O//2. The deposition rate as a function of O//2 flow rate, rf power, and total pressure was obtained. The film composition and properties for samples deposited with different O//2 flow rates were measured. Si-rich films were obtained for the deposition with O//2:SiH//4 flow rate ratio of less than 3. 5. The refractive index increased and film thickness decreased after samples were annealed. Etch rate in a solution containing ten parts of ammonium fluoride and one part hydrofluoric acid (10:1 BHF) and refractive index decreased by increasing the Si:O ratio in samples. The SiH//n bonds were found for all Si-rich films. Both Si-OH and H-OH bonds were observed for all samples.
This review summarizes a variety of beam damage phenomena relating to oxides in (scanning) transmission electron microscopes, and underlines the shortcomings of currently popular mechanisms. These phenomena include mass loss, valence state reduction, phase decomposition, precipitation, gas bubble formation, phase transformation, amorphization and crystallization. Moreover, beam damage is also dependent on specimen thickness, specimen orientation, beam voltage, beam current density and beam size. This article incorporates all of these damage phenomena and experimental dependences into a general description, interpreted by a unified mechanism of damage by induced electric field. The induced electric field is produced by positive charges, which are generated from excitation and ionization. The distribution of the induced electric fields inside a specimen is beam-illumination- and specimen-shape- dependent, and associated with the experimental dependence of beam damage. Broadly speaking, the mechanism operates differently in two types of material. In type I, damage increases the resistivity of the irradiated materials, and is thus divergent, resulting in phase separation. In type II, damage reduces the resistivity of the irradiated materials, and is thus convergent, resulting in phase transformation. Damage by this mechanism is dependent on electron-beam current density. The two experimental thresholds are current density and irradiation time. The mechanism comes into effect when these thresholds are exceeded, below which the conventional mechanisms of knock-on and radiolysis still dominate.
Auger electron spectroscopy (AES) depth profiling was applied for determination of the thickness of a macroscopic size graphene sheet grown on 2 in. 6H-SiC (0 0 0 1) by sublimation epitaxy. The measured depth profile deviated from the expected exponential form showing the presence of an additional, buffer layer. The measured depth profile was compared to the simulated one which allowed the derivation of the thicknesses of the graphene and buffer layers and the Si concentration of buffer layer. It has been shown that the graphene-like buffer layer contains about 30% unsaturated Si. The depth profiling was carried out in several points (diameter 50 μm), which permitted the constructing of a thickness distribution characterizing the uniformity of the graphene sheet.
In this study we compare the thicknesses and optical properties of atomic layer deposited (ALD) Al2O3 films measured using table top and mapping ellipsometry as well as X-ray and optical reflectometry. The thickness of the films is varied in the range of 1-50 nm. ALD samples are used as references with well-controlled composition and thickness, as well as with a good lateral homogeneity. The homogeneity is checked using mapping ellipsometry. Optical models of increasing complexity were developed to take into account both the top (surface roughness on the nanometer scale) and bottom interfaces (buried silicon oxide and interface roughness). The best ellipsometric model was the one using a single interface roughness layer. Since the techniques applied in this work do not measure in vacuum, organic surface contamination even in the sub-nanometer thickness range may cause an offset in the measured layer thicknesses that result in significant systematic errors. The amount of surface contamination is estimated by in situ reflectometry measurement during removal by UV radiation. Taking into account the surface contamination the total thicknesses determined by the different methods were consistent. The linearity of the total thickness with the number of atomic layer deposition cycles was good, with an offset of 1.5 nm, which is in good agreement with the sum of thicknesses of the interface layer, surface nanoroughness, and contamination layer.
The reactions of atomic oxygen with the (100) and (111) surfaces of silicon have been investigated by employing supersonic molecular beam techniques, X-ray photoelectron spectroscopy (XPS), and low-energy ion scattering spectroscopy (ISS). Atomic oxygen adsorbs with unit probability on the clean silicon surface, independent of substrate temperature (110–800 K) and incident mean translational energy (3–16 kcal mol−1). Oxidation of clean silicon with an oxygen atom beam is characterized by wo stages: a “fast” stage that corresponds to oxygen chemisorption in the topmost 2–3 silicon layers; and a “slow” stage that corresponds to oxygen incorporation and oxide film growth. The chemisorption stage is described by first-order Langmuirian kinetics with an apparent saturation coverage of approximately 4 ML O(a), the oxide growth stage by lddirect” logarithmic kinetics, where dx/dt = α exp(−x/L), where x is the oxide thickness. Observation of significant oxidation at substrate temperatures of 110 K suggests that oxie growth in the slow stage may occur bya field-assisted mechanism, where an internal electric field aids transport of oxygen to the underlying silicon substrate layers. XPS and ISS results support a two-dimensional layer-by-layer growth mechanism for oxidation at substrate temperatures below 900 K. At higher temperatures, T ≥ 1050 K, oxide growth is three-dimensional involving nucleation and growth of bulk-like oxide islands even for mean coverages as low as 3 ML O(a). ISS results lend support to the formation of “on-top” adsorbed oxygen atoms that cap silicon dangling bonds at the oxide/gas interface. Coincident bombardment of the silicon substrate with an Ar− ion beam leads to an enhanced rate of oxidation. The enhancement can be understood in terms of a model involving secondary implantation of adsorbed oxygen atoms, coupled with the simultaneous formation of reactive sites (e.g., dangling bonds, vacancies) for oxygen chemisorption. The effect of coincident ion bombardment is reduced at elevated substrate temperatures (∼ 800 K), since the resulting increased propensity for adlayer rearrangement leads to a decrease in the number of active sites for oxygen chemisorption.