Structural and electrical studies of ultrathin layers with Si0.7Ge0.3 nanocrystalsconfined in a SiGe/SiO2 superlatticeE. M. F. Vieira, J. Martín-Sánchez, A. G. Rolo, A. Parisini, M. Buljan, I. Capan, E. Alves, N. P. Barradas, O.

Conde, S. Bernstorff, A. Chahboun, S. Levichev, and M. J. M. Gomes Citation: Journal of Applied Physics 111, 104323 (2012); doi: 10.1063/1.4722278 View online: http://dx.doi.org/10.1063/1.4722278 View Table of Contents: http://scitation.aip.org/content/aip/journal/jap/111/10?ver=pdfcov Published by the AIP Publishing

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

Structural and electrical studies of ultrathin layers with Si0.7Ge0.3

nanocrystals confined in a SiGe/SiO2 superlattice

E. M. F. Vieira,1 J. Martın-Sanchez,1 A. G. Rolo,1 A. Parisini,2 M. Buljan,3 I. Capan,3

E. Alves,4 N. P. Barradas,4 O. Conde,5 S. Bernstorff,6 A. Chahboun,1,7 S. Levichev,1

and M. J. M. Gomes1

1Centre of Physics and Physics Department, University of Minho, 4710 – 057 Braga, Portugal2CNR-IMM Sezione di Bologna, via P. Gobetti 101, 40129 Bologna, Italy3Rud-jer Boskovic Institute, Bijenicka cesta 54, 10000 Zagreb, Croatia4ITN, Ion Beam Laboratory, Unit of Physics and Accelerators, E.N. 10, 2686-953 Sacavem, Portugal5Physics Department and ICEMS, University of Lisbon, 1749-016 Lisboa, Portugal6Sincrotrone Trieste, 34149 Basovizza, Italy7Physics Department, FST Tanger, Tanger, Morocco

(Received 21 December 2011; accepted 26 April 2012; published online 31 May 2012)

In this work, SiGe/SiO2 multi-layer (ML) films with layer thickness in the range of a few

nanometers were successfully fabricated by conventional RF-magnetron sputtering at 350 �C. The

influence of the annealing treatment on SiGe nanocrystals (NCs) formation and crystalline

properties were investigated by Raman spectroscopy and grazing incidence x-ray diffraction. At

the annealing temperature of 800 �C, where well defined SiGe NCs were observed, a thorough

structural investigation of the whole ML structure has been undertaken by Rutherford

backscattering spectroscopy, grazing incidence small angle x-ray scattering, high resolution

transmission electron microscopy, and annular dark field scanning transmission electron

microscopy. Our results show that the onset of local modifications to the ML composition takes

place at this temperature for annealing times of the order of a few tens of minutes with the

formation of defective regions in the upper portion of the ML structure. Only the very first layers

over the Si substrate appear immune to this problem. This finding has been exploited for the

fabrication of a defect free metal-oxide-semiconductor structure with a well-defined single layer

of SiGe NCs. A memory effect attributed to the presence of the SiGe NCs has been demonstrated

by high frequency capacitance-voltage measurements. VC 2012 American Institute of Physics.

[http://dx.doi.org/10.1063/1.4722278]

I. INTRODUCTION

Semiconductor crystalline nanostructures (nanocrystals,

NCs) have been intensively studied during the last two deca-

des due to their unique size dependent physical properties.

Si, Ge, and SiGe NCs have received great attention because

of their non-toxicity, abundance in earth resources, and low

cost maintaining the key advantages of state-of-the-art sili-

con processing, which makes them suitable for applications

in electronics, optoelectronics, and solar cells.1–10 In particu-

lar, SiGe alloys have attracted much interest since the early

1970 s due to the possibility of band-gap engineering by

varying the stoichiometry of the alloy and the excellent Si

and Ge miscibility.11 SiO2 shows incomparable advantages

with respect to other dielectric materials in terms of Si tech-

nology compatibility, and it has also been demonstrated to

be an effective barrier to Ge out-diffusion.12 In the last dec-

ade, most of the effort has been devoted to the production of

Si NCs embedded in SiO2 dielectric matrix by several

techniques.13–16 Concerning SiGe material, NCs production

embedded in SiO2 matrix has been studied by molecular

beam epitaxy (MBE),17 low-pressure chemical vapor deposi-

tion (LPCVD),18 chemical etching methods,19 atom beam

sputtering,20 and rf-magnetron co-sputtering.21,22

For practical applications, it is crucial to control the

NCs morphological and structural properties, such as their

size, shape, density, spatial distribution, and stoichiometry,

which are very important for modelling electronic properties

and optoelectronic devices performance. In particular, non-

volatile memory (NVM) devices based on NCs as discrete

charge storage nodes were first proposed and demonstrated

by Tiwari et al.23 as an excellent alternative to conventional

continuous floating gates NVM. In this regard, SiGe NCs

based NVMs have been successfully demonstrated.4,24 For

an optimal performance, it would be highly desirable to de-

velop a fabrication process that allows one to obtain a high

density of NCs with high size uniformity in a very thin and

well defined two-dimensional layer embedded in a dielectric

material matrix where the inter-distance between the NCs

and the substrate can be maintained constant in large areas.

Usually, in order to obtain self-assembled SiGe NCs embed-

ded in an oxide matrix, a high temperature annealing process

follows the co-deposition of a SiGe-rich matrix oxide layer.

It is well known that high temperature treatments can lead to

undesired Ge diffusion or evaporation throughout the barrier

matrix material with the consequent loss of abrupt and well

defined interfaces.25 Although SiGe NCs in thin layers have

been obtained after annealing of an initially well defined

amorphous SiGe layer deposited by LPCVD in SiO2 matrix

material,26,27 the produced NCs layer structures showed a

clear degradation after performing the annealing process.

0021-8979/2012/111(10)/104323/9/$30.00 VC 2012 American Institute of Physics111, 104323-1

JOURNAL OF APPLIED PHYSICS 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

In this work, we have optimized a fabrication process for

the production of [SiGe/SiO2] multi-layer (ML) structure

grown at 350 �C by RF-magnetron sputtering technique on Si

(100) substrates. The investigation shows an overall ML stabil-

ity with well-organized SiGe NCs confined in 5-nm-thick

layers, after performing a rapid thermal annealing (RTA) treat-

ment at 800 �C for 30 min, in nitrogen atmosphere. However,

for this annealing treatment, the onset of a diffusive phenom-

enon leading to local modifications of the ML composition is

also reported and discussed. We show that as this process does

not affect the very first layers over the Si substrate, metal-

oxide-semiconductor (MOS) structures consisting of a single

SiGe NCs layer confined between SiO2 layers may be success-

fully fabricated in these conditions. Finally, capacitance-

voltage (C-V) measurements performed on MOS capacitor

structures show good charging and discharging capabilities,

demonstrating the suitability of the whole fabrication process

presented for NVM applications.

II. EXPERIMENTAL

A. Samples preparation

The samples presented in this work were prepared

at 350 �C substrate temperature on p-type low resistivity

(1-5 X cm) Si (100) substrates using a commercial Alcatel

SCM650 RF-magnetron sputtering machine. In order to

optimize and study the SiGe NCs production process in thin

layers, we have initially grown an amorphous 20-period

SiGe/SiO2 ML structure by depositing alternated SiGe and

SiO2 thin layers using a composite target of Si (99.999%

purity) plate covered with Ge (99.999% purity) polycrystal-

line pieces and a single SiO2 target, respectively. The ratio

between the Ge pieces and the Si target was 1:4. The

growth rate was approximately 6.7 nm/min and 4.6 nm/min

for SiGe and SiO2 layers, respectively. The layer thick-

nesses were controlled by the deposition time. Before depo-

sition, the native oxide and possible contaminants present

on the substrate surface were thermally desorbed by heating

the Si substrate up to 500 �C for 4 h with a base pressure

(Pbase) of 3� 10�6 mbar in the deposition chamber. The

deposition process was started after introducing an Ar gas

flux in the chamber (Pbase¼ 4� 10�3 mbar) using a radio-

frequency power of 80 W. Finally, the samples were

annealed under inert nitrogen atmosphere (6 mbar) at dif-

ferent temperatures (500–800 �C) and annealing times

(15 and 30 min) to study the crystallization process of the

as-grown amorphous SiGe layers. The annealing was per-

formed using a homemade system (RTA setup) that was

built with three commercially available infrared heaters,

Model 5306 StripIRVR

. Each 5306 StripIR infrared heater is

a 1000 W lamp and polished aluminium reflector heating

system that provides even heat distribution across a 1.5 in.

wide and 5 in. long strip. In order to evaluate the SiGe NCs

layer charging and discharging properties, [Au gate con-

tact/SiO2 control oxide layer/SiGe NCs/SiO2 tunnel oxide

layer/p-Si substrate/Au back-side contact] MOS structures

were fabricated. The gold contacts were deposited by

thermal evaporation at room temperature using a mask with

0.8 mm2 circular openings.

B. Experimental characterizations

Grazing incidence x-ray diffraction (GIXRD) and

Raman scattering measurements were employed to provide

information about the crystallinity, size, and chemical com-

position of the NCs. GIXRD was carried out in a Bruker

AXS D5000 Diffractometer employing CuKa radiation

(k¼ 0.154 nm) at a fixed grazing incidence angle of 1. The

data were collected in the 15�–50� 2h range with a step size

of 0.04� and a step time of 25 s. Raman scattering spectra of

the samples were recorded in backscattering geometry, using

a Jobin-Yvon T64000 system with an optical microanalysis

system and a CCD detector at room temperature using the

488 nm Arþ laser excitation line. The laser beam was

focused on the sample surface with a beam spot size of 1 lm

and a power of 0.5 mW to avoid the heating of the sample.

Silicon TO mode at 521 cm�1 was used as calibration fre-

quency reference. Rutherford backscattering (RBS) and elas-

tic recoil detection analysis were used for the investigation

of the compositional depth profile along the ML structure in

the growth direction. For RBS measurements, a 2.0 MeV4Heþ beam was employed with tilt angles of 78� and 82� in

order to obtain sufficient resolution to resolve the ML struc-

ture depth profile on the nanometre scale. Annular dark field

scanning transmission electron microscopy (ADF-STEM),

high resolution transmission electron microscopy (HRTEM),

and selected area electron diffraction (SAED) were

employed to investigate the SiGe/SiO2 ML structure, the

quality of the interfaces between the layers, and the nature of

the SiGe NCs. To this end, a FEI Tecnai F20 transmission

electron microscope, operating at 200 kV, was used. Samples

for TEM observations were prepared in cross-sectional ori-

entation. The ML structure of samples was also evaluated by

grazing incidence small angle x-ray scattering (GISAXS) at

the SAXS beamline of the Elettra Synchrotron (Trieste,

Italy), using a photon energy of 8 keV. Two-dimensional

(2D) GISAXS maps were obtained at grazing incidence

angles slightly above the critical angle for total external

reflection. For the electrical characterization, high-frequency

(1 MHz) C-V measurements were recorded at room tempera-

ture with a SULA Technologies spectrometer.

III. RESULTS AND DISCUSSION

GIXRD patterns and Raman spectra for as-grown and

annealed samples are shown on the Figs. 1(a) and 1(b),

respectively. The samples were annealed at temperatures

(Tann) ranging from 500 �C to 800 �C for an annealing time

(tann) of 15 min. In the insets to these figures, corresponding

spectra obtained on a sample annealed at 800 �C for 30 min

are also reported. As shown by GIXRD (Fig. 1(a)), no evi-

dence of crystalline features is observed for the as-grown

sample and up to an annealing temperature of Tann¼ 500 �C.

However, for Tann� 600 �C, SiGe NCs formation is observed

through the definition of broad peaks centered at diffraction

angles falling in between the {111} and {220} values of the

Si and Ge cubic diamond structures.28

The results obtained by GIXRD are essentially confirmed

by the Raman investigation reported in Fig. 1(b). For

Tann� 600 �C, three Raman peaks around 300, 400, and

104323-2 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

500 cm�1 are observed corresponding to Ge-Ge, Ge-Si, and

Si-Si vibrations modes of a crystalline SiGe alloy, respec-

tively.29 The Raman peaks present a high asymmetrically

shape and are red shifted with respect to the bulk material val-

ues. For all the samples, the Si-Si peak is more intense than

the Ge-Ge peak, indicating that our films are Si-rich alloy.29

In agreement with the GIXRD results, for Tann < 600 �C, only

broad features are observed, which corresponds to the amor-

phous phase. For comparison, the Raman signal coming from

the substrate is plotted in Fig. 1(b) showing a peak corre-

sponding to bulk crystalline Si (521 cm�1). From these obser-

vations, we conclude that the SiGe layers crystallization

temperature is about 700 �C where the GIXRD and Raman

peaks features corresponding to a SiGe crystalline alloy are

observed. In thick amorphous films, the crystallization is a ho-

mogeneous process throughout the film that can be explained

by classical nucleation theory.30 However, an exponential

increase of the crystallization temperature with decreasing

layer thickness has been reported for ultrathin Si/SiO2

(Ref. 31) and Ge/SiO2 (Ref. 32) amorphous superlattices. The

value we obtained for our Si1-xGex/SiO2 ML is well inside the

crystallization temperature range that may be estimated, at a

comparable value of the layer thickness, from these two limit-

ing cases, i.e., 750 �C (Ref. 31) and 600 �C,32 respectively.

The previously reported results demonstrate that crystalline

features start to appear at a temperature of 700 �C and well

defined nano-crystalline SiGe structures are observed at

Tann¼ 800 �C, thus in what follows, only the latter annealing

temperature will be considered.

The GIXRD experimental results obtained at 800 �C(Fig. 1(a)) were also used to determine the average size of

the SiGe NCs. After performing a Lorentzian curve fitting to

the experimental diffraction peaks, from their position and

full width at half maximum (FWHM), an average SiGe NCs

size of about 4.6 6 0.2 nm was calculated by using the

Debye–Scherrer formula.33 Within the experimental error,

no clear NCs size variation was found for the different

annealing times (15 and 30 min). A Ge content (x) value of

0.32 in the Si1�xGex alloy was calculated from the linear

Vegard law,34 by the equation:

aSiGe ¼ xaGe þ ð1� xÞaSi; (1)

where aSiGe is determined from the experimental data, and

aGe¼ 0.5658 nm and aSi¼ 0.5431 nm are the Ge and Si bulk

lattice parameters, respectively.

In order to investigate the ML structure evolution after

the longest treatment (Tann¼ 800 �C and tann¼ 30 min), RBS

(scattering angle of 160�) and ERDA (recoil angle of 24�)measurements were performed on as-grown and annealed

samples. Fig. 2 shows the experimental RBS spectra

obtained at an 82� tilt grazing angle on a ML sample before

(a) and after the RTA treatment (c). In this figure, the results

of least-squares fits of a realistic physical model of the sam-

ple and of the RBS parameters to the experimental data,

based on the NDF software,35 are also reported. The ERDA

data, from where we determined the hydrogen concentration,

are not shown here. It should be noted that, from the experi-

mental RBS spectra, only about four Ge-rich layers, at the

top of the ML structure, are clearly resolved. Average values

and uncertainties for the layer thickness and composition

were obtained from these data and assumed to be representa-

tive of the whole ML structure. The total thickness of the

ML, as measured by RBS, is consistent with the average val-

ues determined for the top layers, confirming the validity of

this assumption. Furthermore, significant deviations in the

thickness of individual deeper layers would lead to measura-

ble effects in the RBS data. The resulting fitted depth con-

centration profiles are shown in Figs. 2(b) and 2(d) for the

as-grown and annealed samples, respectively. As expected,

these profiles show a regular sequence of SiGe/SiO2 layers

in both cases. Interestingly, even after performing a high

temperature and long annealing treatment, the ML structure

is preserved within the RBS measurements accuracy. The

multilayer oscillations seem to be slightly smaller after

annealing, but this is difficult to quantify, given the limited

FIG. 1. (a) GIXRD patterns of as-grown and annealed samples from 500 �Cto 800 �C for 15 min. The vertical dashed lines show the position of Si and

Ge diffraction peaks. (b) Corresponding Raman spectrum of each sample:

Ge-Ge, Ge-Si, and Si-Si optical vibration modes are shown. Insets: same for

30 min annealing time.

104323-3 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

statistics in the small oscillations. No significant change in

all the elements concentration (Si, Ge, and O) along the ML

structure was observed after performing the annealing pro-

cess. From the RBS depth profile data assuming the bulk

density of SiO2 and a weighted average of the atomic den-

sities of Si and Ge for SiGe, the average thickness values for

the SiGe and SiO2 layers, after the annealing treatment, were

found to be 5.62 6 0.5 nm and 5.80 6 0.5 nm, respectively,

with similar values obtained for the as-grown sample.

GISAXS analysis of the films was performed in order to

obtain more insights into the structural properties of the ML,

namely on the quality of the interfaces between the layers.

We would like to point out that the main contribution to the

scattering signal in these experiments is coming from the

upper layers of the ML. GISAXS maps obtained for the as-

grown and annealed films are shown in Fig. 3(a). Both maps

have very similar intensity distributions, showing that the

main contribution to the scattered intensity comes from the

interfaces between the layers.36 The SiGe NCs are sur-

rounded by amorphous SiO2 and other SiGe NCs, so they

cannot be well resolved by the GISAXS method. This hin-

ders the possibility to obtain an accurate determination of

their structural properties and size. However, from the very

strong Bragg sheet intensities that arise from correlated inter-

face roughness,36 it is possible to determine the dependence

of the layer thicknesses and of the interface roughnesses on

the annealing treatment. With this aim, the profiles of the

Bragg sheet along the lines indicated in Fig. 3(a) have been

fitted. The fitting was performed by using standard formulas

for correlated roughness calculated in the distorted-wave

Born approximation (DWBA).36,37 The fitting parameters

were thicknesses of the layers; surface roughness; and lateral

and vertical correlation lengths. The extracted 1D experi-

mental profiles and the corresponding fittings are shown in

Figs. 3(b) and 3(c), respectively. The extracted profiles are

very similar for the as-grown and annealed films. Small shift

of the Bragg peak position to larger Qz values in Fig. 3(c) is

caused by slightly different incidence angles of the probing

beam. Actually, after correction for the refraction effects

(included in the fit), the positions of all Bragg peaks are prac-

tically the same. The incidence angles were chosen to ensure

maximal intensity of the signal; different values indicate

changes in the refraction index of the material with anneal-

ing. The results of the fitting are summarized in Table I,

where it is shown that no significant variation of the film

thicknesses is seen after the annealing treatment. Moreover,

a slight decrease of the interface roughness as well as a small

increase in the lateral correlation length is found upon

annealing. The vertical correlation length was approximately

constant during the annealing.

The above reported RBS and GISAXS experimental

results have given a rather complete picture of the as-grown

FIG. 2. RBS spectra of the as-grown (SiGe/SiO2) ML sample before (a) and after (c) an RTA annealing at 800 �C for 30 min. Concentration depth profiles,

resulting from a least squares fitting of a realistic physical model of the sample and of the RBS parameters to the experimental data (red lines in (a) and (c)),

are shown in (b) and (d), in the case of the as-grown and the annealed sample, respectively.

104323-4 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

SiGe/SiO2 ML and have suggested that the annealing treat-

ment affects only marginally its structure. However, two

main reasons suggest the need of a further structural investi-

gation based on a more direct and local characterization

technique. On the one hand, as already noted, the previously

employed techniques have probed essentially the top portion

of the ML, a direct observation of the whole ML structure

being still lacking. On the other hand, the above results do

correspond to large area averages that may overlook the pos-

sible presence of structural defects or local fluctuations of

the layer composition. Hence, cross-sectional TEM and

STEM observations were performed on the as-grown and

annealed ML samples.

In Figs. 4(a) and 4(b), the results of a X-sectional ADF-

STEM investigation of the whole SiGe/SiO2 ML structure,

before and after an 800 �C annealing treatment for 30 min,

are reported. As shown in these images, in both the cases,

throughout the whole ML structure, the SiGe layers appear

wavy with the exception of the first bottom layers starting

from the substrate. However, the extent of this effect is

clearly different in the two cases, layer roughnesses with

greater amplitudes and shorter lateral correlation lengths,

being clearly observed after the annealing treatment, Fig.

4(b). A number of information on the ML structure can be

gained by an analysis of the diffraction patterns reported in

Figs. 4(c) and 4(d) obtained starting from the images in Figs.

4(a) and 4(b), respectively. In these patterns, the periodicity

of a perfect ML structure is coded in a series of equispaced

diffraction spots aligned in a direction parallel to the surface

normal. From the spacings of the diffraction features

observed in this direction in Figs. 4(c) and 4(d), it results

that, within the accuracy of the measurements, the ML perio-

dicity does not change after annealing; the values obtained

being 11.5 6 0.3 nm and 11.4 6 0.3 nm, respectively. This is

in agreement with what previously found by both RBS and

GISAXS measurements extending the validity of this finding

to the whole ML structure. However, in Figs. 4(c) and 4(d), a

pronounced arcing of the diffraction spots is observed. This

effect reflects the occurrence of correlated ML roughnesses

and its extension their amplitude. Namely, as particularly

evident after the annealing treatment (Fig. 4(d)), the fact that

this arcing is not circular but appears to follow an elliptical

shape indicates the existence of a local strain in a direction

normal to these undulated regions, the period of these

deformed regions being shorter than that observed along the

surface normal. The observed increase in the interface rough-

ness has a local character. Therefore, TEM and STEM are

more suited to study this phenomenon than GISAXS, the lat-

ter technique giving essentially statistical data over large

areas as mentioned above. This justifies the apparent discrep-

ancy between both techniques (see Figures 3 and 4). Degra-

dation in the structural quality of ML structures fabricated

FIG. 3. GISAXS maps of the as-grown

film and the film annealed at 800 �C for

30 min (a). 1D profiles were extracted

along the yellow dashed lines for the fit-

ting procedure. The extracted 1D profiles

(black circles) and the corresponding fits

(red lines) for the as-grown and annealed

films (b). (c) The grazing incidence

angles were 0.28�, 0.26�, and 0.24� for

as-grown, 15 min, and 30 min annealed

films, respectively.

TABLE I. Structural properties of SiGe/SiO2 interfaces obtained by

GISAXS analysis. ML period (T), interface roughness (r), lateral (g), and

vertical (w) correlation lengths are shown. All values are given in nm.

Sample T r g W

As-grown 11.5 6 0.2 2.0 6 0.2 10 6 2 24 6 2

Tann¼ 800 �C;

tann¼ 15 min 11.6 6 0.2 1.9 6 0.2 10 6 2 20 6 2

Tann¼ 800 �C;

tann¼ 30 min 11.9 6 0.2 1.8 6 0.2 12 6 2 20 6 2

FIG. 4. Cross-sectional ADF-STEM images of the whole ML SiGe/SiO2

structure before (a) and after (b) a 800 �C thermal annealing for 30 min. In

(c) and (d), the corresponding diffractograms are reported.

104323-5 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

by magnetron-sputter deposition is well known.38 This phe-

nomenon is apparently similar to that observed in our case, a

small intrinsic roughness in individual layers of a ML system

leads to a large cumulative roughness toward the top of the

ML. The same trend is observed for the lateral roughness

along the ML that depending on cumulative effects is

enhanced on the top portion of the ML. Theoretical calcula-

tions made by Payne et al.39 have shown that this interfacial

correlation increases with the square root of the layer index

(the lowest index corresponding to the layer closest to the

substrate). In other words, flatter and sharper interfaces are

expected for the first layers, whereas due to cumulative

effects of interface roughness and its correlation along the

ML, the top layers are expected to have rougher interfaces.

The origin of these phenomena has been attributed to a

too low surface energy flux during sputter deposition that

reducing atomic rearrangement gives rise to the observed

interface roughness accumulation.38 However, in our case, it

is clear that, in the as-grown sample, the ML structure,

although not perfect, is by far more regular than that

observed after thermal annealing. This suggests that an addi-

tional phenomenon is taking place during the latter process.

In Fig. 5, the results of an ADF-STEM and EDS x-ray inves-

tigation performed at higher magnification, are reported. In

the ADF-STEM image of the top portion of the ML reported

in Fig. 5(a), two arrows are aligned along an almost vertical

line where a faint layer contrast is observed. This elongated

region may be seen more clearly in Fig. 5(b), where the

same image shown in Fig. 5(a) is displayed in a false color

scale. Thus, these Z-contrast images suggest that a consider-

able modification of the ML composition has taken place in

this region, up to the onset of a local ML structure disrup-

tion. This is at variance with what observed in the as-grown

samples where layer continuity appears to be maintained.

Moreover, the fact that these type of defects, present in the

annealed samples, are visible in images representing the pro-

jected ML structure (Figs. 5(a) and 5(b)) also indicates that

they should possess a considerable extension along the direc-

tion of observation, i.e., the normal to these cross-sectional

images. A clear demonstration that a modification of the ML

composition has taken place after the annealing process is

shown in Fig. 5(c) where EDS x-ray profiles taken along the

AB line marked in Fig. 5(b), are reported. In this plot, the

absolute values of the composition in atomic % should be

taken with care as, owing to the signal integration over the

TEM sample thickness, roughness superposition could give

rise to a possible mixing of signals coming from adjacent

layers. However, in Fig. 5(c), at the defect site, an unambigu-

ous increase in oxygen content and a parallel Ge depletion

are observed. It is worth noting that, in a previous work,26

high temperature long furnace annealing (up to 1 h) of low

pressure chemical vapor deposition (LPCVD) SiGe/SiO2

FIG. 5. ADF-STEM and EDS x-ray

investigation of the ML sample annealed

at 800 �C for 30 min. (a) ADF-STEM

image of the top portion of the ML; (b)

the image in (a) is represented in a false

color scale; (c) EDS x-ray profiles taken

along the line AB marked in (b).

104323-6 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

MLs resulted in ML composition degradation owing to a Ge

diffusion towards the Si substrate. In our case, ADF-STEM

analysis of the bottom region of the ML (not shown here)

showed that the first layers from the Si substrate are essen-

tially unaffected by the annealing treatment. Compositional

variations start to appear from the third or fourth layer from

the substrate and increase towards the top of the ML, follow-

ing the same cumulative trend followed by the layer rough-

ness. This suggests a possible link between the two

phenomena. As a matter of fact, these defects are observed at

the edge of the layers undulations, Figs. 4(b), 5(a), and 5(b),

which as previously noted represents regions with different

strains. At present, the precise mechanism responsible of the

enhancement of the layer roughness observed after the

annealing treatment is still a matter of speculation. Here, we

just bound ourselves to note that SiGe heterogeneous nuclea-

tion at the upper and lower layer interfaces and compressive

strain, known to lead to an enhancement of the Ge self-

diffusion in Si1-x Gex layers,40 could play an important role

in the observed phenomena. Finally, it is important to remind

here that the very first layer over the Si substrate appears

immune to these problems as this fact will be exploited for

the fabrication of a MOS structure to be reported in the

following.

In Fig. 6, HRTEM micrographs of the top and bottom

regions of the SiGe/SiO2 ML structure are reported. SiGe

NCs are clearly observed only within the dark contrasted

SiGe layers as marked in the HRTEM image with circles.

These layers appear formed by 3 to 5 nm sized SiGe crystal-

line grains. In the inset to Fig. 6(b), SAED pattern obtained

on a larger area than that visible in Fig. 6 but approximately

centered on the same TEM sample region is shown. The

observed diffraction rings demonstrate the presence of NCs

showing no preferred growth orientation. After careful cali-

bration of the SAED pattern, an average {111} interplanar

distance of 0.319 6 0.006 nm may be measured. This value

is in keeping with the previously reported GIXRD estimate,

confirming the existence of a Si-rich SiGe alloy. Finally, as

shown in Fig. 6, HRTEM observations also reveal the pres-

ence of local variations of the NCs lattice fringe spacings.

These distortions in the 5 nm-thick SiGe layers are the result

of a crystalline growth constrained between the embedding

6 nm-thick SiO2 layers. Interestingly, a good quality of the Si

substrate/SiO2 interface is observed with no clearly visible

defects.

Finally, the structural and electrical properties of [Au gate

contact/SiO2 control oxide layer/SiGe NCs/SiO2 tunnel oxide

layer/p-Si substrate/Au back-side contact] MOS structure were

studied. For SiGe NCs formation, an annealing process at

Tann¼ 800 �C for 15 min was performed. A cross-sectional

HRTEM image of this structure is shown in Fig. 7(a) where ab-

rupt SiGe/SiO2 and SiO2/Si interfaces are observed. The pres-

ence of SiGe NCs with {111} interplanar distance of

0.318 6 0.004 nm is clearly observed in agreement with the

results previously obtained on the ML sample by GIXRD and

TEM. In particular, this value is between those corresponding

to pure Si (0.314 nm) and Ge (0.327 nm) for the {111} family

crystal planes. Fig. 7(b) shows averaged intensity line profiles

centred on the line A-B marked in Fig. 7(a). Besides the profile

obtained on the original micrograph, in black in Fig. 7(b), a

profile obtained on the very same image region after the appli-

cation of a low-pass filter in Fourier space to the image is also

reported in red. In this figure, the interfaces are determined by

the positions of the averaged profile inflection points.41,42

The filtering procedure improves the accuracy of this determi-

nation. The resulting thicknesses of the SiO2 control oxide,

SiGe NCs layer, and SiO2 tunnel oxide are 8.3 6 0.2 nm,

8.2 6 0.3 nm, and 5.6 6 0.2 nm, respectively. The absence of

SiGe lattice fringes in some region of the SiGe layer can be

FIG. 6. Cross-sectional HRTEM micrographs of a SiGe/SiO2 ML sample

annealed at 800 �C for 30 min. In (a) and (b), regions close to the sample sur-

face and the Si substrate are shown, respectively. 5 nm-thick SiGe layers

with nanocrystals (region circled with blue dotted lines) separated by SiO2

layers are shown. The SAED diffraction pattern, in the inset shown in (b),

obtained on a larger area, confirms the crystallinity of the SiGe NCs. Scale

markers in the HREM and SAED micrographs correspond to 5 nm and

2 nm�1, respectively.

FIG. 7. (a) Cross-sectional HRTEM

micrograph of a MOS structure com-

posed by a SiGe NCs layer confined

between two SiO2 layers on a Si sub-

strate after an annealing at 800 �C for

15 min. NCs are clearly identified by the

presence of {111} SiGe lattice fringes.

In (b), averaged intensity line profiles

centered on the line A–B shown in (a)

are reported. Black and red profiles are

obtained on the original and low-pass fil-

tered (allowed spatial frequencies up to

1.7 nm�1) micrograph, respectively.

104323-7 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

explained by assuming strong disorientation of NCs with

respect to the Si substrate or/and by the presence of amorphous

pockets between the SiGe NCs. From these observations, it is

not clear if isolated NCs are obtained in the layer.

Corresponding high-frequency (1 MHz) C-V curves for

a sweeping voltage of 6 3 V are presented in Fig. 8. In order

to avoid dynamic recharging during measurement, a ramp

rate of 0.05 V/s has been used. For comparison, the C-V

curve measured on a sample, containing a completely amor-

phous SiGe layer, is shown in the inset to this figure. A

counter-clockwise hysteresis loop in the C-V curve with flat-

band voltage shift value of DVFB¼ 0.46 V is only obtained

when SiGe NCs are present in the layer. The voltage sweep

was made from inversion to accumulation regimes and back.

It is well known that the origin of a C-V hysteresis loop

can be attributed to the presence of traps localized: (i) inside

NCs or at their interfaces with the dielectric matrix (NCs-

related traps); (ii) in the dielectric matrix; and (iii) at the

interface between the dielectric matrix and the Si substrate.

In this regard, given that such hysteresis is not observed for

the sample where the SiGe layer is amorphous, we attribute

the charge trapping mainly to the presence of SiGe NCs.

Therefore, this result demonstrates that the stacked layers do

not act as charge storage centres. As mentioned before, the

possibility that the SiGe layer is completely crystalline with

NCs forming the grains of a continuous nanocrystalline layer

may not be ruled out. In this case, we cannot ignore the lat-

eral charge loss that could take place in the layer because of

the absence of clear and evident isolated and discrete storage

centres, which could justify the relative small memory win-

dow observed in the hysteresis loop. However, this could be

counteracted by growing thinner SiGe layers where isolated

NCs could be obtained as we plan to investigate in a future

work. A more detailed study about the charge trapping

mechanism is also needed and will be done in another work.

IV. CONCLUSIONS

We have demonstrated a fabrication process of SiGe

NCs/SiO2 ML structures with nanometer layer thicknesses

by RF-magnetron sputtering. The optimization of the SiGe

NCs formation process was carried-out by Raman spectros-

copy and GIXRD measurements. This investigation has

shown that for annealing times of a few tens of minutes,

SiGe crystalline structures start to appear at a temperature of

700 �C, and well defined SiGe NCs are observed at 800 �C.

However, a thorough characterization of the effects of a

RTA annealing at the latter temperature on the overall ML

structure, undertaken by RBS, GISAXS, HRTEM, and ADF-

STEM investigations, has shown that, for annealing times of

the order of a few tens of minutes, the initial stage of a diffu-

sive phenomenon giving rise to local modifications to the

ML composition occurs via the formation of defective

regions in the upper portion of the ML. This process does

not affect the first SiGe/SiO2 layers over the Si substrate;

hence, MOS structures with a single SiGe NCs layer con-

fined between SiO2 layers obtained as described in this work

can be successfully fabricated. High frequency C-V meas-

urements on the MOS structure have shown the presence of

a hysteresis phenomenon in the C-V curve that is interpreted

in terms of charge trapping by the SiGe NCs layer. This find-

ing indicates that this fabrication process is suitable for

memory applications.

ACKNOWLEDGMENTS

This study has been partially funded by: (i) FEDER

funds through the COMPETE program “Programa Operacio-

nal Factores de Competitividade” and by Portuguese funds

through Portuguese Foundation for Science and Technology

(FCT) in the frame of the Project PTDC/FIS/70194/2006;

(ii) the transnational access framework of the ANNA Eu Pro-

ject (Contract No. 026134 RII3) through the funding of the

ANNA_TA_UC9_RP006 proposal; (iii) ELETTRA Syn-

chrotron Radiation Center for the measurements at the

SAXS beamline funding received from the European Com-

munity’s Seventh Framework Programme (FP7/2007–2013)

under Grant Agreement No. 226716; (iv) Scientific and

Technological Cooperation Program between Portugal

(FCT) and Morocco (CNRST)-2010/2011; (v) European

COST MP0901-NanoTP Action. E.M.F.V., J.M.S., and S.L.

are grateful for financial support through the FCT Grant Nos.

SFRH/BD/45410/2008, SFRH/BPD/64850/2009, and SFRH/

BPD/26532/2006, respectively. M.B. acknowledges support

from the Ministry of Science Education and Sports, Republic

Croatia (Project No. 098-0982886-2866). I.C. acknowledges

support from the Unity through Knowledge Fund. The

authors would like also to thank Jose Santos for technical

support.

1E. Kapetanakis, P. Normand, D. Tsoukalas, K. Beltsios, J. Stoemenos, S.

Zhang, and J. van den Berg, Appl. Phys. Lett. 77, 3450 (2000).2S. Gardelis, P. Manousiadis, and A. Nassiopoulou, Nanoscale Res. Lett. 6,

227 (2011).3M. Yang, T. P. Chen, Z. Liu, J. I. Wong, W. L. Zhang, S. Zhang, and Y.

Liu, J. Appl. Phys. 106, 103701 (2009).4Y. Liu, S. Tang, and S. K. Banerjee, Appl. Phys. Lett. 88, 213504 (2006).

FIG. 8. High-frequency (1 MHz) capacitance-voltage curves of the annealed

MOS structure (a). For comparison, a C-V curve of an amorphous SiGe

layer is presented in the inset.

104323-8 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00

5L. Khomenkova, N. Korsunska, V. Yukhimchuk, B. Jumayev, T. Torchynska,

A. V. Hernandez, A. Many, Y. Goldstein, E. Savir, and J. Jedrzejewski,

J. Lumin. 102–103, 705 (2003).6X. J. Hao, A. P. Podhorodecki, Y. S. Shen, G. Zatryb, J. Misiewicz, and

M. A. Green, Nanotechnology 20, 485703 (2009).7M. H. Liao, C. Y. Yu, T. H. Guo, C. H. Lin, and C. W. Liu, IEEE Electron

Device Lett. 27, 252 (2006).8S. Takeoka, K. Toshikiyo, M. Fujii, S. Hayashi, and K. Yamamoto, Phys.

Rev. B 61, 15988 (2000).9T. Kirchartz, K. Seino, J.-M. Wagner, U. Rau, and F. Bechstedt, J. Appl.

Phys. 105, 104511 (2009).10S. K. Kim, C. H. Cho, B. H. Kim, S. J. Park, and J. W. Lee, Appl. Phys.

Lett. 95, 143120 (2009).11G. Abstreiter, Light Emission from Silicon: From Physics to Devices,

Semiconductors and Semimetals Vol. 49, edited by D. J. Lockwood

(Academic, San Diego, 1998), p. 38, 39.12H. G. Chew, W. K. Choi, Y. L. Foo, F. Zheng, W. K. Chim, Z. J. Voon, K. C.

Seow, E. A. Fitzgerald, and D. M. Y. Lai, Nanotechnology 17, 1964 (2006).13L. Patrone, D. Nelson, V. I. Safarov, M. Sentis, W. Marine, and S. Giorgio,

J. Appl. Phys. 87, 3829 (2000).14O. M. Feroughi, C. Sternemann, Ch. J. Sahle, M. A. Schroer, H.

Sternemann, H. Conrad, A. Hohl, G. T. Seidler, J. Bradley, T. T.

Fister, M. Balasubramanian, A. Sakko, K. Pirkkalainen, K. Hamalainen,

and M. Tolan, Appl. Phys. Lett. 96, 081912 (2010).15X. Y. Chen, Y. F. Lu, L. J. Tang, Y. H. Wu, B. J. Cho, X. J. Xu, J. R.

Dong, and W. D. Song, J. Appl. Phys. 97, 014910 (2005).16X. W. Du, L. Y. Liu, P. Yao, and L. Cui, J. Appl. Phys. 100, 076102

(2006).17I. N. Demchenko, K. Lawniczak-Jablonska, S. Kret, A. V. Novikov, J. Y.

Laval, M. Zak, A. Szczepanska, A. N. Yabloskiy, and J. F. Krasilnik,

Nanotechnology 18, 115711 (2007).18M. Avella, A. C. Prieto, J. Jimenez, A. Rodrıguez, J. Sangrador, T.

Rodrıguez, M. I. Ortiz, and C. Ballesteros, Mater. Sci. Eng., B 147, 200

(2008).19A. Malachias, S. Kycia, G. Medeiros-Ribeiro, R. Magalhaes-Paniago, T. I.

Kamins, and R. S. Williams, Phys. Rev. Lett. 91, 176101 (2003).20K. U. Joshi, D. Kabiraj, A. M. Narsale, D. K. Avasthi, T. K. Gundurao, T.

N. Warang, and D. C. Kothari, Surf. Coat. Technol. 203, 2497 (2009).21N. A. P. Mogaddam, A. S. Alagoz, S. Yerci, R. Turan, S. Foss, and T. G.

Finstad, J. Appl. Phys. 104, 124306 (2008).

22M. Buljan, S. R. C. Pinto, R. J. Kashtiban, A. G. Rolo, A. Chahboun, U.

Bangert, S. Levichev, V. Holy, and M. J. M. Gomes, J. Appl. Phys. 106,

084319 (2009).23S. Tiwari, F. Rana, H. Hanafi, A. Hartstein, E. F. Crabbe, and K. Chan,

Appl. Phys. Lett. 68, 1377 (1996).24K. Dong-Won, F. E. Prins, K. Taehoon, H. Sungbo, C. Lee, K. Dim-Lee,

and S. K. Banerjee, IEEE Trans. Electron Devices 50, 510 (2003).25V. Beyer, J. Borany, and M. Klimenkov, J. Appl. Phys. 101, 094507

(2007).26A. Kling, A. Rodrıguez, J. Sangrador, M. I. Ortiz, T. Rodrıguez, C.

Ballesteros, and J. C. Soares, Nucl. Instrum. Methods Phys. Res. B 266,

1397 (2008).27A. Rodrıguez, T. Rodrıguez, A. Prieto, J. Jimenez, A. Kling, C. Ballesteros,

and J. Sangrador, J. Electron. Mater. 39, 1194 (2010).28C. W. Hwang, M. K. Ryu, K. B. Kim, S. C. Lee, and C. S. Kim, J. Appl.

Phys. 77, 3042 (1995).29M. I. Alonso and K. Winer, Phys. Rev. B 39, 10056 (1989).30C. Spinella, S. Lombardo, and F. Priolo, J. Appl. Phys. 84, 5383 (1998).31M. Zacharias, J. Blasing, P. Veit, L. Tsybeskov, K. Hirschman, and P. M.

Fauchet, Appl. Phys. Lett. 74, 2614 (1999).32G. V. M. Williams, A. Bittar, and H. J. Trodah, J. Appl. Phys. 67, 1874

(1990).33J. I. Langford and A. J. C. Wilson, J. Appl. Crystallogr. 11, 102 (1978).34A. R. Denton and N. W. Ashcroft, Phys. Rev. A 43, 3161 (1991).35N. P. Barradas, C. Jeynes, and M. A. Harry, Nucl. Instrum. Methods Phys.

Res. B 136–138, 1163 (1998).36U. Pietsch, V. Holy, and T. Baumbach, High-Resolution X-Ray Scattering:

From Thin Films to Lateral Nanostructures, 2nd ed. (Springer-Verlag,

New York, 2004).37V. Holy, C. Giannini, L. Tapfer, T. Marschner, and W. Stolz. Phys. Rev. B

55, 9960 (1997).38S. P. Vernon, D. G. Stearns, and R. S. Rosen, Appl. Opt. 32, 6969

(1993).39A. P. Payne and B. M. Clemens, Phys. Rev. B 47, 2289 (1993).40N. R. Zangenberg, J. Lundsgaard Hansen, J. Fage-Pedersen, and A.

Nylandsted Larsen, Phys. Rev. Lett. 87, 125901 (2001).41M. Ferri, S. Solmi, A. Parisini, M. Bersani, D. Giubertoni, and M. Barozzi,

J. Appl. Phys. 99, 113508 (2006).42A. Parisini, V. Morandi, S. Solmi, P. G. Merli, D. Giubertoni, M. Bersani,

and J. A. van den Berg, Appl. Phys. Lett. 92, 261907 (2008).

104323-9 Vieira et al. J. Appl. Phys. 111, 104323 (2012)

[This article is copyrighted as indicated in the article. Reuse of AIP content is subject to the terms at: http://scitation.aip.org/termsconditions. Downloaded to ] IP:

89.155.9.77 On: Mon, 30 Dec 2013 23:41:00