Thin Solid Films 534 (2013) 90–99
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Thin Solid Films
journal homepage: www.elsevier.com/locate/tsf
Growth of crystalline silicon nanowires on nickel-coated silicon wafer beneath
sputtered amorphous carbon
Feng Ji Li a, Sam Zhang a,⁎, Jun Hua Kong b, Jun Guo b, Xue Bo Cao b, Bo Li c
a
b
c
School of Mechanical and Aerospace Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singapore
School of Materials Science and Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singapore
Central Iron and Steel Research institute, 76 South Xueyuanlu Rd, Haidan District, Beijing, 100081, China
a r t i c l e
i n f o
Article history:
Received 21 March 2012
Received in revised form 1 February 2013
Accepted 1 February 2013
Available online 16 February 2013
Keywords:
Rapid thermal annealing
Crystalline silicon
Nanowires
Nickel catalyst
Oxidation
Solid–liquid–solid growth
a b s t r a c t
Growth of crystalline silicon nanowire of controllable diameter directly from Si wafer opens up another avenue for its application in solar cells and optical sensing. Crystalline Si nanowire can be directly grown
from Si wafer upon rapid thermal annealing in the presence of the catalyst such as nickel (Ni). However,
the accompanying oxidation immediately changes the crystalline Si nanowire to amorphous SiOx. In this
study, amorphous carbon layer was sputtered to on top of the catalyst Ni layer to retard the oxidation. Scanning electron microscope, transmission electron microscope, Raman spectroscopy and X-ray photoelectron
spectroscopy were employed to characterize the wires and oxidation process. A model was developed to explain the growth and oxidation process of the crystalline Si nanowire.
© 2013 Elsevier B.V. All rights reserved.
1. Introduction
Following the discovery of carbon nanotubes [1], one-dimensional
(1-D) anisotropic inorganic structures with nanometer diameters,
such as nanowires, nanotubes and nanorods have been stimulating
great research interest and becoming the hottest research subjects
in recent years [2,3]. As 1-D nanostructure, Si-based nanostructures play
critical roles in interconnects and functional components in mesoscopic
electronic, optical devices and energy storage [4]. Their mechanical, electronic and optical properties strongly rely on dimensionality and lattice
orientation [2]. For instance, the ultimate mechanical strength of silica
nanowire is in excess of 12GPa and increases with the decrease of the diameter [5]. Diamond lattice Si only has an indirect 1.1-eV band gap showing weak infrared (IR) emission, whereas linear trans-polysilane chains
exhibit a 3.89-eV direct band gap showing strong ultraviolet (UV) emission [6]. Silica nanowires could emit strong blue [7,8], red [7], green [6]
and even UV light [7,8]. Moreover, it is capable of guiding light within
the visible and near infrared spectral ranges with low optical losses [9].
However, the synthesis of Si-based nanostructures with desired
dimensionality or morphologies is technically challenging [10]. Generally speaking, top-down and bottom-up approaches are the main routes
to synthesize Si nanowires. Top-down method prepares Si nanowires
via dimensional reduction of bulk Si by lithography and etching,
which is usually time-consuming, expensive and difficult to achieve
⁎ Corresponding author. Tel.: +65 67904400.
E-mail address: [email protected] (S. Zhang).
0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.tsf.2013.02.007
uniform diameters below 10 nm [11,12]. Bottom-up approach is an assembly process joining Si atoms often via a vapor–liquid–solid (VLS)
mechanism in generating single crystalline Si nanowires with diameters ranging from 5 nm to several hundreds of nanometers [3]. In the
mechanism [13,14], a liquid eutectic metal (such as Au, Ni, Fe, etc.) or
alloy droplet composed of metal catalyst components and nanowire
materials (such as Si, C etc.) is formed under the reaction conditions [10].
This liquid droplet serves as a preferred site for absorption of gas-phase
reactant. Once supersaturated, it becomes the nucleation site for crystallization. During the growth, the catalyst droplet directs the nanowire's
growth orientation and defines the diameter of the crystalline nanowire.
Ultimately, the growth terminates when the temperature is below the
eutectic temperature of the alloy or the reactant is no longer available.
The formation of eutectic metal/Si liquid droplets and the establishment
of the symmetry-breaking solid–liquid interface is the key for the
one-dimensional crystal growth. However, it has relatively low yield of
production, and the use of flammable or toxic precursor gases such as
SiH4/H2, SiCl4 or Si oxide vapor are the disadvantages. So far, chemical
vapor deposition [15], laser ablation [3], thermal evaporation [16–18],
and thermal annealing have been applied to synthesize Si-based 1-D
nanostructures [10,19–21]. For instance, Goodey et al. synthesized Si
nanowires from SiH4 at 500 °C by VLS mechanism using Au as the catalyst [22]. Alfredo et al. synthesized single-crystalline Si/SiOx nanowires
with crystalline cores vary from 6 to 20 nm, and length of 1 to 30 μm
via combining laser ablation nanometer-diameter iron catalyst cluster
formation [3]. Justin et al. obtained defect-free Si nanowires with uniform diameters ranging from 4 to 5 nm, length of several microns via a
F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
supercritical fluid solution-phase approach with Au nanocrystal as the
catalyst [2].
To avoid using flammable or toxic precursor gas as the Si source in
VLS, solid–liquid–solid (SLS) growth of amorphous Si, amorphous SiOx
or crystalline-Si/amorphous-SiOx (i.e. c-Si/a-SiOx) nanowires directly
uses Si wafer as the source and adopts Ni [20], Au [23,24], NiO [18], or
Fe [25] as catalysts under nitrogen or argon atmosphere. Park et al. synthesized c-Si/a-SiOx nanowires by heating a NiO-catalyzed Si substrate
under a reductive environment of WO3, C and Ar atmosphere via
SLS growth [18]. A tentative SLS mechanism for the growth of the
nanowire based on chemical reaction has been proposed: CO gas
from the carbothermal reaction of WO3/C reacted with SiOx to produce SiO and CO2, where SiO2 is formed from Si–Ni–O liquid droplets
in the NiO/Si interface. Si and SiOx are generated through the disproportionation of SiO in the film [18]. The c-Si/a-SiOx nanowire is grown via
the phase separation of Si and SiOx from the super-saturated Si–Ni–O
droplets by the cooling effects of Ar flow. Taking consideration of the
oxidation of the droplet at the beginning stage of the growth, Nie et
al. developed a solid–liquid–solid growth mechanism catalyzed by the
temperature-dependent alloy, metallic and ionic Fe (Fe–FeOx) to describe the growth of Si–SiOx core–shell nanowires during directly
heating Fe-coated Si wafers at three different temperatures. In the
mechanism, the formation of 1-D amorphous SiOx nanowires was
grown from the fully oxidized Fe–Si alloying droplets as a result of
the continuously feeding Si into the alloying droplets, whereas partly oxidized large ones form Fe–Si alloying cores and Fe–O–Si shells,
which result in Si–SiOx core–shell nanowires.
Except Nie et al., almost all the other mechanisms overlooked the
serious oxidation by residual oxygen in the reaction chamber at high
temperature, since the reaction is often under an inert gas atmosphere
rather than a high vacuum environment. Moreover, if these reactions
take place, Si crystal might be embedded in amorphous SiO2 matrix in
the form of film, rather than Si crystalline core encapsulated in SiOx
sheath in the form of the nanowire [26]. Therefore, the tentative mechanisms (reaction plus phase separation, or disproportionation reaction)
are challenged to explain why the as-grown Si nanostructure is in the
form of crystalline Si core wrapped in amorphous SiOx sheath. Amorphous SiOx nanowires can be grown directly by heating a NiO-catalyzed
silicon substrate without using WO3/graphite powder in a reductive environment in a tube furnace [18], or by thermal heating Au or Ni thin
film-coated Si substrates [27]. The residual oxygen is unavoidable in the
chamber. The formation of amorphous all-SiOx nanowires is attributed
to the serious oxidation of as-grown Si wires. Therefore, in order to
grow crystalline Si nanowires, the oxidation of as-grown Si nanowires
must be retarded. In summary, the main challenges in Si-related nanostructures are: (1) difficulty in controlling the dimensionality and morphologies. (2) Incredulity of the proposed SLS mechanisms based on
phase separation of Si and SiOx or disproportionation reaction of SiO
vapor in explaining the growth of c-Si/a-SiOx nanowires. (3) Inability in
explaining the formation of the amorphous SiOx nanowires by the disproportionation reaction mechanism.
As a promising alternative to standard furnace annealing, rapid thermal annealing presents advantages in precise control of annealing temperature and the annealing profile, ambient purity, short annealing
times (from 1 s to 3 min), cycle time, and process flexibility. The fundamental flexibility in creating different types of thermal processes arises
from the dynamic control of the heat source temperature, which permits
fast heating combined with dynamic optimization of temperature uniformity [28,29]. However, the annealing chamber normally is under a
gas atmosphere instead of a vacuum environment due to the high cost.
Alternatives of retarding the unavoidable oxidation must be worked
out to grow crystalline Si nanowires. Our previous work reported the
growth of amorphous all-SiOx nanowires via 3 min rapid thermal
annealing of 345-nm a-C/Ni film using a-C layer as the oxidation retardation layer [30]. However, c-Si/a-SiOx nanowires were still not
observed due to the serious oxidation. This paper sputtered thicker
91
carbon film on a thin layer of Ni using Si wafer as substrate, followed
by rapid thermal annealing for different duration of time to control
the growth of the c-Si/a-SiOx nanowires. It is shown that the structure
of c-Si/a-SiOx nanowires is controlled by the thickness of the a-C layer
and the annealing duration. In the end, an oxidation-accompanying
SLS growth model is developed to explain the formation of the c-Si
nanowires, amorphous all-SiOx nanowires and their transition.
2. Experimental details
2.1. Deposition of the film
The a-C/Ni film was prepared in an E303A magnetron sputtering
system (Penta-Vacuum, Singapore). N-type Si (100) wafer was used
as the substrate (10 mm × 10 mm in area, 475 μm in thickness and
0.5 nm in Rq). Before loading into the sputtering chamber, the substrate was ultrasonically cleaned in acetone for 20 min, followed by
10 min in alcohol. Once the base pressure reached 1.0 × 10 −3 Pa, Ar
gas was introduced at a flow rate of 50 standard cubic centimeters per
minute (sccm). High energy Ar+ ions were generated at a substrate
bias of −300 V for substrate cleaning for 15 min to further clean the surface contamination. Afterwards, a 5-min sputtering of Ni target (99.99%
in purity) in 0.47 Pa pressure at room temperature resulted in about
30-nm thick Ni catalyst layer on the substrate. Si and C were subsequently co-sputtered from Si and graphite target (99.999% in purity) on the Ni
layer under the same pressure. Amorphous-C films of around 130 nm
and 570 nm in thickness were respectively co-sputtered for 15 min and
60 min. As a comparison study, a single Ni film was also deposited from
Ni target for 5 min. The sputtering power densities for Ni, Si and C were
respectively RF 1.85 W/cm2, RF 2.47 W/cm2 and DC 12.22 W/cm2.
2.2. Rapid thermal annealing
The as-sputtered films underwent rapid thermal annealing (RTA,
Jipelec Jetfirst 100 rapid thermal processor) in an Ar ambient at 1100 °C
for 60 s and 180 s, respectively. The RTA chamber was purged 10 times
with Ar (99.999% in purity) at 2000 sccm before ramping up to 600 °C
at 58 °C/s. After dwelling for 6 s at 600 °C, the chamber was further
ramped up to 1100 °C at 41.7 °C/s, and then held for 60 s and 180 s, respectively. After annealing, the chamber was firstly cooled down to
500 °C at 43 °C/s, followed by natural cooling to room temperature. During the whole annealing process (ramping, holding, and cooling), the inflow of the Ar gas was kept at 2000 sccm to maintain chamber pressure of
around 1.08×105 Pa.
2.3. Characterization
The surface chemical states of the films were studied by X-ray
photoelectron spectroscopy (XPS, Kratos) using a Kratos AXIS spectrometer with monochromatic Al Kα (1486.71 eV) radiation (10 mA
and 15 kV). The electron energy analyzer was operated at 0.1 eV scanning step size on a low magnification scan spot with 220 μm in diameter. The atomic concentration of each element is determined by the
ratio of the corresponding fitted core-level peak area over the sum of
the fitted peak area of all the elements. The respective core-level XPS
curve fitting was performed after a Shirley background subtraction by
nonlinear least square fitting using a mixed Gauss/Lorentz function
after charge-correcting by positioning C 1s peak to 284.8 eV. Field emission scanning electron microscopy (JEOL, JBM-7600F, JEOL Ltd., Japan,
5-kV operating voltage) showed that wire-like nanostructures were
formed on the surface of the as-annealed film surface. High resolution transmission electron microscope (HRTEM, JEOL 2010 and 2100F,
200-kV operating voltage) revealed that the wires were of a crystalline
core enveloped by an amorphous sheath. HRTEM samples were prepared by using a sharp stainless tweezers to scratch the as-grown
wires off the substrate surface into a small plastic container of acetone,
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F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
and subsequently followed by 1 min ultrasonic agitation for dispersion.
The size of the wire was calculated by randomly taking the mean value
of 10 wires. Energy dispersive X-ray spectroscopy (EDX, EDAX Inc., US)
cross-sectional elemental mapping and selected area electron diffraction
(SAED) confirmed that the core was crystalline Si and the encapsulation
was amorphous SiOx. The carbon sp2 bonding structure was examined
with a Renishaw Raman spectroscope at a wavelength of 633 nm radiation excited with a He–Ne laser.
3. Results and discussion
3.1. Surface chemical state of the film before and after annealing
Fig. 1(a) and (b) shows the XPS survey scan spectra of the
as-sputtered and the 3-min annealed 570-nm a-C film. Note that no
signal for Ni 2p (binding energy ranging from 885 to 847.9 eV) is
detected in the spectrum except for Si 2p, Si 2s, O 1s, O 2s and C 1s,
which means that Ni still locates at the bottom of the film. As seen,
upon annealing, the signal of C 1s is dramatically trailed off, while
the intensities of O 1s, O 2s, Si 2s and Si 2p are greatly boosted up. Before annealing, the C 1s content (79.34 at.%) is ultrahigh in comparison with Si 2p (4.16 at.%) and O 1s (16.50 at.%). The amount of Si in
the film is very small. After annealing, there is a tremendous decrease
in C 1s content (5.2 at.%) and sharp increase in O 1s (67.10 at.%) and
Si 2p (27.69 at.%).
To evaluate the chemical state of silicon in the as-sputtered film
and the oxidation of silicon after annealing, the Si 2p core-level peaks
of the film before and after annealing are illustrated in Fig. 7(c) and
(d). The five oxidation states of Sin+ (n= 0, 1, 2, 3 and 4) corresponding
to the five chemical structures of Si, Si2O, SiO, Si2O3 and SiO2, respectively [31], exist in nonstoichiometric SiOx (0b x b 2) films. The position of
the five synthetic peaks of Si 2p is determined thus: 99.6 eV in Si0,
100.6 eV in Si1+, 101.6 eV in Si2+, 102.6 eV in Si3+, and 103.6 eV in
Si 4+ [32,33]. The full width at half maximum of each peak is fixed to
be 1.90 eV during fitting. After charge correcting the Si 2p peak by positioning the major component in C 1s peak to 284.8 eV, the above peak
positions are used to fit the Si 2p core-level lines. A variation in the position of each peak of ±0.1 eV is considered acceptable. With the
possible existence of the five oxidation states in mind, after subtracting
an integrated Shirley function background, the curve fitting is conducted
by decomposing the spectrum into the Si 2p1/2 and Si 2p3/2 partner lines
for the 5 oxidization states without any pre-adjudication. Although the
spin-orbit splitting and the Si 2p1/2 and Si 2p3/2 intensity ratio are not
fixed in the fitting, the fitting yields reasonable results. The fitting yields
that the spin-orbit splitting is 0.6±0.05 eV and the intensity ratio is
equal to the statistical value of 1:2 for all of the oxidation states.
As seen in Fig. 1(c) and (d), prolonging the annealing to 3 min, the
atomic concentration of the Si4+ component is up-shifted and the
atomic concentration of sub-oxide components is decreased. Note that
Si0 and Si1+ peaks don't appear in the whole process indicating Si 0
does not exist in the a-C film. Huge difference of the atomic concentrations of the other four Si environments is that: for the as-sputtered film,
the contents of Si3+ (38.73 at.%) and Si2+ (57.26 at.%) are high in comparison with Si4+ (4.01 at.%); for the annealed one, there is a tremendous increase in the content of Si4+ (59.79 at.%) and sharp decrease
in the concentration of Si3+ (29.53 at.%) and Si2+ (10.68 at.%). Upon
annealing, the film is transformed from a random bonding model structure into distinct mixtures of graphitic carbon and SiO2-like. The reasons
are summarized as follows. The annealing system is not a high vacuum
device thus the presence of residual O2 is unavoidable. In addition, even
if the chamber pressure with Ar (1.08× 105 Pa) is a little higher than the
outside atmosphere (1.01 × 105 Pa), a small amount of O2 could still be
introduced into the chamber. Therefore, the additional O2 might have
come from the leakage of and the residual O2 in the annealing system.
The presence of the additional O2 and the high annealing temperature
result in the removal of carbon and corresponding increase of oxidization
in the film. This, taken into consideration with an increase in the Si 2p
binding energy, intimates the formation of oxidized coatings which are
more SiOx>1.
In the previous work [30], the thickness of the a-C film is around
345 nm deposited under the same target power density as in the
present study. Since XPS can only detect the chemical composition
and bonding structures beneath the surface to around 10 nm. Therefore, the XPS survey spectra and high resolution spectra in [30] and
the present work look similar. However, there are still some differences
as listed below. a) The chemical composition shows difference: In [30],
Fig. 1. XPS spectra of the (a) as-sputtered film, (b) 3-min annealed films; deconvolution of Si 2p core-level spectra of the (c) as-sputtered film, and (d) 3-min annealed film.
F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
93
Fig. 2. Cross section of the (a) 570-nm a-C/30-nm Ni film, (b) 130-nm a-C/30-nm Ni film, and (c) 50-nm Ni film. Scale bar is 100 nm.
in as-sputtered film, C 1s (90.04±1.20 at.%), Si 2p (5.21±0.40 at.%) and
O 1s (4.75±0.97 at.%), in as-annealed film, C1s (18.33±4.58 at.%), Si 2p
(47.46 ± 3.49 at.%) and O 1s (34.09 ± 2.22 at.%). b) The sub-oxides
are different: In [30], the spectrum doesn't have Si 0 and Si 1 + for
the as-sputtered film, the Si sub-oxides in the as-annealed film was
completely converted to Si 4 +.
To summarize, Ni catalyst particle doesn't move along with the
wire. It remains at the bottom, rather than at the tip of the wire. The
amount of Si in the as-deposited film was very small. The as-annealed
film has a close chemical composition to stoichiometric SiO2. Si0 doesn't
exist in the as-sputtered film.
3.2. Morphology, chemistry and structure
3.2.1. Cross section of the film
The typical cross section of the a-C film of 570 nm and 130 nm in
thickness are respectively shown in Fig. 2(a) and (b). The Ni-catalyst
layer of around 30 nm is clearly observed between the a-C top layer
and Si substrate in Fig. 2(b). However, it is not very clearly seen in
Fig. 2(a) due to the low contrast between the thick top layer and the Ni
layer. But the white color of the Ni layer is also seen as highlighted by
the arrow at the interface. The thickness of the single Ni film is around
50 nm as clear shown in Fig. 1(c). The details of the as-sputtered
a-C/Ni and Ni film are summarized in Table 1.
3.2.2. Morphology of the wire
Fig. 3(a) shows the wires on the surface of the 600-nm a-C/Ni film
after 3-min annealing. The wires are of around 70 nm in diameter, up
to 10 μm in length, entangling with each other. The wires become inhomogeneous with bright and transparent outer sheaths and a white
inner core indicating a heterogeneous structure. Fig. 3(b) shows bundles of wires growing from the cracked film. Fig. 3(c) shows the wires
grown after 3-min annealing of the 160-nm a-C/Ni film. The wires are
of homogeneous structure and of round ends, lying on the film surface entangling into each other. Fig. 3(d) shows the wires grown via
3-min annealing of the single Ni layer without the a-C layer on the
top. In summary, wire-like nanostructures grow during annealing of
the a-C/Ni/Si-wafer and Ni/Si-wafer samples. Ni particle doesn't appear at the tip of the wire.
Fig. 4(a) and (b) shows the structure of the nanowires after 1-min
annealing of the 600-nm a-C/Ni film. The wires have a mean outer diameter of 57.23±11.80 nm, a mean core diameter of 20.97±5.07 nm and a
mean sheath thickness of 18.00±2.90 nm. The ratio of the core diameter
over the sheath thickness is 1.16±0.14. After prolonging the annealing to
3 min, the wire's outer diameter increases to 84.75±25.53 nm, and the
core diameter reduces a lot, down to 8.62±0.45 nm and the sheath
thickness increased to 36.49±14.41 nm (cf. Fig. 4(c) and (d)). The
ratio of the core diameter over the sheath thickness is 0.25±0.09,
which is dramatically decreased compared to the 1-min annealing. The
end of the core is surrounded by a 12.1 nm thin sheath layer as shown
in the inset of Fig. 4(d). Obviously, oxidation takes place inward to consume the core. Fig. 4(e) and (f) shows the nanowires via 3-min annealing
of the 160-nm a-C/Ni film. The wire in Fig. 4(f) is of 70.05 nm in diameter
showing a homogeneous amorphous structure without an inside core.
Fig. 4(g) and (h) shows the structure of the wires after 3-min annealing
of the 50-nm Ni film without the a-C film. Homogeneous amorphous
wires are also observed without the black core. The dark lines randomly
entangled with the Si nanowires do not belong to Si nanowires. They
are the residual carbon dissolved in the solution during the preparation
of the TEM samples. The details of the as-grown nanowires are summarized in Table 1.
In summary, with or without the a-C film on the top of the Ni
layer, homogeneous or heterogeneous nanowires are able to form. It
also confirms that Si in the wire is from Si wafer. However, the cover
layer of a-C does provide oxidation reduction. As shown in Table 1,
within the same annealing duration of 3 min, the ratio of the core diameter over the sheath thickness decreases from 0.25± 0.09 to 0 as the a-C
Table 1
Details of the as-deposited a-C/Ni, Ni film and the as-grown nanowires.
Film structure
Thickness of
the a-C layer
(nm)
Thickness of
the Ni layer
(nm)
Annealing
duration
(min)
Ratio of core diameter
over sheath thickness
a-C/Ni
a-C/Ni
a-C/Ni
Ni
570
570
130
0
30
30
30
50
1
3
3
3
1.16 ± 0.14
0.25 ± 0.09
0
0
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F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
Fig. 3. Wires grown via 3-min annealing of the (a)–(b) 600-nm a-C/Ni film, white arrow indicating the center of the wire, (c) 160-nm a-C/Ni film, and (d) 50-nm Ni film.
layer thickness reduces from 570 nm to 130 nm (cf. Fig. 4(c)–(f)). Besides, for the same a-C layer thickness of 570 nm, prolonging the
annealing from 1 min to 3 min gives rise to reduced ratio of the core
diameter over the sheath thickness from 1.16 ± 0.14 to 0.25 ± 0.09
(cf. Fig. 4(a)–(d)).
3.2.3. Chemistry and structure of the wire
Fig. 5(a) shows the cross-sectional elemental mapping of the
as-grown wire after 1-min annealing of the 600-nm a-C/Ni film. As
seen, it consists of Si and O only. The intensity of Si and O exhibits almost the same parabolic function along the cross-sectional route, not
Fig. 4. TEM images of as-grown wires via (a)–(b) 1-min annealing of the 600-nm a-C/Ni film, (c)–(d) 3-min annealing of the 600-nm a-C/Ni film, (e)–(f) 3-min annealing of the
160-nm a-C/Ni film, and (g)–(h) 3-min annealing of the 50-nm Ni film.
F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
only indicating stable atomic concentration ratio (1:2) between Si
and O along the route, but also implying that the cross section is
very round. Particularly, the intensity of O has a small plateau while
that of the Si continues to increase while approaching the core of
the nanowire, which could be regarded as an implication of the Si
core.
Fig. 5(b) shows a single Si nanowire consisting of a dark core with
a diameter of 12.31 nm surrounded by a gray sheath of 21.17 nm in
thickness. The ratio of the core diameter over the sheath thickness
is 0.58. The nanowire is straight with uniform smooth morphology
along the growth direction. Inset in Fig. 5(b) shows the corresponding
SAED patterns of the selected circle area of the wire. A unit cell of
cubic structure (lattice parameter = 5.43 Å) was consistent with the
diffraction pattern. Indexing of the SAED patterns shows the [111]
zone axis of single-crystalline Si, which suggests
h that
i the nanowire is
constructed with crystalline Si growing along 110 orientation [34],
surrounded by amorphous SiOx sheath (i.e. c-Si/a-SiOx). Fig. 5(c) shows
the high resolution TEM image of the defect-free c-Si/a-SiOx nanowire
via 1-min annealing of the 570-nm-thick a-C film. The wire is consisted
of a crystalline Si core with diameter of 9.50 nm surrounded by an amorphous SiOx sheath with thickness of 11.29nm. The lattice spacing of
3.84 Å agrees closely with that of the 110 lattice structure of silicon.
95
There is a dihedral angel of 120° between 110 and 011 (lattice
spacing, 3.84 Å) in the atom-resolved TEM image. The 2D fast Fourier transform (FFT) pattern of the lattice resolved core image is
shown in the inset of Fig. 5(c). It is from the typical Si [111] zone axis.
The FFT pattern in conjunctionh with
i the lattice-resolved core image reveals the growth is along the 011 orientation.
3.2.4. Carbon bonding structure
Fig. 6 shows the transformation of carbon sp 2 bonding structure
after 3-min annealing of the 570-nm film. The Raman spectra are
deconvoluted into D band (at ~ 1350 cm −1 A1g mode) and G band
(at ~1580 cm−1, E2g mode) using a mixture of Lorentzian and Gaussian
functions [35,36]. D peak is the breathing mode of those sp 2 sites only in
rings [35]. Its intensity is directly dependent on the presence of the
six-fold rings. G peak is attributed to the in-plane stretching vibration
of any pair of sp2 bonded carbon atoms, either in C_C chains or in aromatic rings [35]. ID/IG increases as the number of rings per cluster increases and the fraction of chain groups falls [37].
After annealing, G peak is up-shifted from 1533 cm−1 to 1588 cm−1.
The corresponding ID/IG increases from 0.98 to 1.77. The second order
bands at approximate 2643 cm−1 and 2908 cm−1 are extruded. The
band at 2643 cm−1 is the most intensive one and attributed to the first
Fig. 5. (a) EDX cross sectional elemental mapping spectrum characterizes the chemical composition of the nanowire via 1-min annealing of the 600-nm a-C/Ni film to be
constructed with Si core and SiOx (x≈2) sheath, (b) TEM image of nanowire via 1-min annealing of the 600-nm a-C/Ni film, inset shows the SAED pattern of the core, and (c) high resolution TEM image of the nanowire grown via 1-min annealing of the 600-nm a-C/Ni film showing the crystalline core and the amorphous sheath, the insets showing the corresponding
FFT pattern of the core.
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F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
Fig. 6. Carbon sp2 bonding structure of the 600-nm a-C/Ni film (a) as-sputtered and (b)3-min annealed.
overtone of the D band (around 2 times of 1350) [38,39]. An additional
higher order band at 2908 cm−1 has been assigned to a combination
of the G and D modes characteristic for disturbed graphitic structures.
The shape of the Raman spectra before and after annealing looks similar
to the previous work in [30]. That's because they all underwent serious
graphitization and oxidation after heating at 1100 °C. However, the difference still exists after deconvoluting of the spectrum into G band and
D band. The ID/IG ratio in [30] increases from to 0.94 to 1.29, while that
of the present study increases from 0.98 to 1.77, indicating the different
extent of graphitization. To summarize, once the temperature surpasses
a certain point, the graphitization inclination and the crystallization is
dramatically accelerated and strengthened. The tremendous increase of
ID/IG, the huge up-shift of the G-band position, the appearance of the
second-order bands at high Raman shift centers are attributed to growing sp2 clusters and the conversion from sp3 to sp2 bonding configuration during the rapid thermal annealing.
3.3. Growth mechanism
3.3.1. Ni catalyst
Fig. 7(a) shows a spherical Ni particle of around 38.1 nm in diameter embedded at the end of a c-Si/a-SiOx nanowire. The hemispherical
particle is enveloped by a thin SiOx layer (thickness of 5.7 nm) at the periphery. The wire has an outer diameter of 32.8 nm, tapered inner diameter of 14.1 nm at the beginning (indicated by d1) and of 6.0 nm at
the top (indicated by d2). The tapered Si core indicates less oxidation
at the wire end, implying the oxidation retardation effect of a-C layer.
It has a clear interface with the Si core, suggesting that the Ni particle
has an orientation relationship or is partially coherent with the wire.
The spherical shape suggests that the particle was in liquid phase during
growth [13,14]. The typical lattice fringes of the Ni particle are clearly
observed over the entire particle as magnified in Fig. 7(b). Fig. 7(c)
shows a cucurbit-like Ni particle directing the growth of the amorphous
SiOx nanowires. The particle has length of 119.0 nm and width of
75.1 nm, surrounded by a layer of SiOx. The indexed SAED patterns
shown in Fig. 7(d) are from Ni [113] zone axis [40]. These observations
confirm that the Ni–Si phase does exist at the end of the wire and contributes to the growth of the wires.
3.3.2. Si source
VLS growth of Si nanowires includes: a) absorption of Si atoms into
the liquid eutectic metal or alloy droplets from the vapor Si phase, the
liquid droplet serves as a preferential site; b) precipitation of Si nanowires
once super-saturation of Si reaches, the liquid droplet serves as a nucleation site for crystallization; and c) termination of the growth when the
temperature is below the eutectic temperature of the catalyst metal or
alloy, or the reactant is no longer available [20]. However, under the
present conditions, Si concentration in the vapor phase is negligible at
the growth temperature because the specific surface/volume ratio of
bulk Si substrate is extremely low compared with that of the laser
ablation [3].
It is also different from the solid–liquid–solid growth of Park et al.
[18], who grew c-Si/a-SiOx nanowires via 3 h heating a NiO-coated
(solution method) Si substrate under a reductive environment of WO3,
C and Ar atmosphere. The WO3 and C were in the form of mixed powders
besides the NiO-coated Si substrate. However, in the present experiment,
the sputtered carbon top layer was separated by a layer of Ni. The longest
annealing time is only 3 min. In addition, amorphous SiOx nanowires
due to serious oxidation during annealing were also grown without
sputtering the carbon layer on the top of nickel as shown in Fig. 3(d).
Therefore, the disproportionation reaction of SiO vapor was excluded.
The carbothermal reduction reaction between carbon and oxygen should
be considered when the carbon layer is on the top. However, the function
of the reaction is only to retard the oxidation rather than generate Si and
SiO2 through the disproportionation reaction of SiO vapor as in the work
of Park et al. In addition, Si0 is not in the film as confirmed in Fig. 1(c).
Therefore, the effect of small amount of Si in the as-deposited film can
be ignored. Si is neither from the vapor nor from the as-deposited a-C
layer. The only credible source is the bulk Si substrate.
3.3.3. Oxidation-accompanying solid–liquid–solid growth model
The binary Ni–Si diagram shows that the eutectic temperature of
NiSi2 is around 993 °C [41]. Due to the melting effect of small-size
grains [42], the eutectic compound NiSi2 can begin to form at a temperature lower than 964 °C. The sputtered Ni layer is able to melt
on the Si wafer substrate at the annealing temperature of 1100 °C,
and form NiSi2 eutectic droplets. Considering the surface oxidation of
the silicon wafer, the NiSi2 droplet is surrounded by Si–O. Because of
the relatively high solubility of Si in NiSi2 eutectic alloy, more Si atoms
would diffuse through the solid (substrate) liquid interface into the liquid phase. A second liquid–solid interface forms when the liquid phase
becomes supersaturated with Si atoms due to thermal or compositional
fluctuations, resulting in the growth of crystalline Si nanowire. The alloy
droplets encapsulated with SiOx amorphous layer have been confirmed
F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
97
Fig. 7. (a) Ni catalyst particle of 38.1 nm in diameter encapsulated at the end of the c-Si/a-SiOx wire, (b) high resolution TEM image of the Ni catalyst clearly showing the lattice
fringes, (c) cucurbit-like Ni particle encapsulated at the end of the amorphous SiOx nanowire and (d) the SAED patterns of the Ni particle are from the Ni [113] zone axis.
in Fig. 7(a) and (c). The amorphous SiOx sheath is simultaneously
formed due to oxidation. The growth involves not only the evolution
of solid–liquid–solid phases, but also relates to the simultaneous oxidation of the as-grown Si nanowire. Different from the Si-SiOx core–shell
nanowires synthesized by Nie et al. without sputtering the top a-C
layer [25], the oxidation from the residual oxygen is effectively retarded
in this work. The whole structure evolution responsible for the growth
of c-Si/a-SiOx nanowires is thus developed taking place in the following
three steps as schematically sketched in Fig. 8. To simplify the growth
mechanism responsible for the evolution of Si nanowires, only the
annealing duration is directly taken into consideration, whereas the
a-C top layer is not drawn in the schematic diagram.
(I) Initiation of the Si nanowire
1) Formation of the Ni–Si–O droplet. As sketched in Fig. 8(I), Si
atoms are generated from the interface of Ni/Si wafer substrate
at the annealing temperature of 1100 °C (which is higher than
the eutectic temperature of the Ni–Si alloy), and then dissolve
into the liquid Ni to form Ni–Si droplet. The oxygen on the Si
wafer surface partially oxidizes Ni–Si alloy to form Ni–Si–O
on the shell. 2) Precipitation of the c-Si/a-SiOx nanowire. Upon
super-saturation of Si in the Ni–Si droplet, a concentration gradient is formed due to excess Si at the interface between the Ni
layer and Si substrate. Meanwhile, the gradual burning out of
carbon in the film at the Ar/O2 atmosphere gives rise to porous
or cracked a-C film. As a result, the crystalline Si/amorphous
SiOx (c-Si/a-SiOx) nanowire of dIC in core diameter precipitates
upward from the Si-supersaturated Ni–Si–O droplet. The crystalline Si core is depicted with a preferred growth orientation
b110> in the diagram.
(II) Oxidation and growth of the Si nanowire
With the precipitation of c-Si/a-SiOx wire, the residual oxygen
in the reactor results in the inward oxidation of the c-Si core
(dIIC b dIC). Note that the a-C film was partially cracked during
the annealing. A little bit of residual oxygen could further penetrate inside to oxidize the Ni–Si–O droplet.
(III) Transition from Si to SiOx nanowire
As time goes, the c-Si core is completely consumed and it completes the transition from c-Si to a-SiOx nanowire. Fig. 8(III)
shows the completion of the transition from c-Si to amorphous
all-SiOx. SiOx layer envelopes the Ni particle inactivating it and
ending the growth of the wire, leaving a SiOx-rich surface.
High resolution TEM analysis reveals the presence of the Ni–Si–O
phase at the end of the wire, where the Ni crystal is embedded at
the end of the c-Si/a-SiOx nanowire (cf. Fig. 7(a) and (c)), and oxygen
is resulted from the oxidation. Within the same time of annealing, as
a result of the complete oxidation during annealing of the 160-nm
98
F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
Fig. 8. Schematic growth and oxidation of the crystalline Si nanowire: (I) free Si atoms diffuse from Si substrate which dissolves into the Ni atoms until reaching a Ni–Si alloy
super-saturation (Ni–Si eutectic droplets), where Si is expelled from the Ni–Si core as a crystalline nanowire of dIC in diameter, the wire is depicted with a preferred b110> orientation. The oxygen on the Si wafer surface partially oxidizes Ni–Si alloy to form Ni–Si–O on the shell, thus the a-SiOx sheath surrounds the c-Si. The outer SiOx sheath is surrounded
by the residual O2, (II) the c-Si core becomes smaller (dIIC b dIC) due to the inward oxidation, and (III) termination of the growth as the annealing time is up. The entire Si wire becomes amorphous SiOx nanowire as a result of the complete oxidation. The Ni–Si–O droplet becomes a spherical Ni crystal surrounded by SiOx shell.
a-C/Ni film, the entire c-Si core is oxidized into amorphous SiOx
nanowire (cf. Fig. 4(e)–(f)), whereas the 600-nm thick a-C/Ni layer
gives rise to c-Si/a-SiOx nanowires (cf. Fig. 4(c)–(d)). Furthermore,
for the same a-C top layer, shorter annealing duration gives rise to
the c-Si/a-SiOx nanowire of bigger crystalline core diameter, thus higher
ratio of the core diameter over the sheath thickness (cf. Fig. 4(a)–(b)).
The driving force for such a continuous diffusion of Si atoms from
the substrate into the liquid droplets, and then through the Ni–Si liquid
eutectic droplets into solid nanowires could be attributed to the kinetic
and environment aspects. From the kinetic point of view, the concentration gradient and super-saturation due to the fluctuation in Ni–Si droplets are the driving force of the growth. From the environment point,
the carrier Ar gas surrounds the semi-spherical shape Ni–Si droplets
and exchange energy and momentum with the atoms in the Ni–Si droplet, resulting in overcooling to the droplets. Such an overcooling is
critical to initiate the preferential unidirectional growth of c-Si/a-SiOx
nanowires. The resultant amorphous SiOx sheath is attributed to the oxidation reaction between crystalline Si nanowire and residual O2. The a-C
film retards or attenuates the oxidation rate of the solid–liquid–solid
growth of precipitated Si nanowire during rapid thermal annealing.
Plus, the a-C gives rise to porous or cracked film to direct the growth of
Si nanowire during the annealing. Parameters such as the thickness of
the a-C top layer, the annealing duration, Ni layer thickness, Ar flow
rate, and annealing temperature might play important roles on the morphology and structure of the nanowire.
In summary, at the annealing temperature of 1100 °C, eutectic Ni–Si
droplet forms at the Ni/Si wafer interface; here, Si atoms diffuse into the
Ni–Si droplet; upon super-saturation, Si precipitates and forms crystalline Si which serves as the seed for Si nanowire. With time, crystalline Si
wire forms. Meanwhile, residual oxygen in the annealing chamber oxidizes the crystalline Si wire to form amorphous SiOx sheath, giving rise
to c-Si/SiOx structure. Prolonged annealing or reduced a-C layer thickness results in further inward oxidation of the crystalline Si core into
amorphous SiOx till the complete consumption of the c-Si core and
thus the formation of amorphous all-SiOx nanowires.
4. Conclusions
Underneath amorphous carbon (a-C) layer, crystalline silicon
nanowires can be grown directly from nickel-coated Si wafer substrate via rapid thermal annealing. The a-C layer has retarded the
oxidation of the Si nanowire. Without the a-C layer, Si nanowire is
completely oxidized to amorphous SiOx. Within the same annealing duration, thicker layer gives rise to bigger core diameter. For the same
layer thickness, prolonged annealing results in smaller core diameter. An oxidation-accompanying solid–liquid–solid model explains
the growth of the Si nanowire and its oxidation process.
Acknowledgments
This work was supported by the Singapore Ministry of Education's
Research Grant T208A1218 ARC4/08.
F.J. Li et al. / Thin Solid Films 534 (2013) 90–99
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