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Silicon nanocrystals on amorphous silicon carbide alloy thin films: Control of filmproperties and nanocrystals growth
Jérémy Barbé a,b,⁎, Ling Xie c, Klaus Leifer c, Pascal Faucherand a, Christine Morin a, Dario Rapisarda a,Eric De Vito a, Kremena Makasheva b,d, Bernard Despax b,d, Simon Perraud a
a CEA, Liten, 17 rue des Martyrs, 38054 Grenoble Cedex 9, Franceb Université de Toulouse, UPS, INPT, LAPLACE (Laboratoire Plasma et Conversion d'Energie), 118 route de Narbonne, 31062 Toulouse, Francec Department of Engineering Sciences, Uppsala University, Box 534, S-751 21 Uppsala, Swedend CNRS, LAPLACE, F-31062 Toulouse, France
⁎ Corresponding author at: Université de Toulouse, UPPlasma et Conversion d'Energie), 118 route de Narbonn
Article history:Received 11 January 2012Received in revised form 21 August 2012Accepted 24 August 2012Available online 31 August 2012
Keywords:Low pressure chemical vapor depositionSiliconNanocrystalsPlasma enhanced chemical vapor depositionAmorphous silicon carbideTransmission electron microscopyX-ray photoelectron spectroscopySilicon carbide
The present study demonstrates the growth of silicon nanocrystals on amorphous silicon carbide alloy thinfilms. Amorphous silicon carbide films [a-Si1−xCx:H (with xb0.3)] were obtained by plasma enhanced chem-ical vapor deposition from a mixture of silane and methane diluted in hydrogen. The effect of varying the pre-cursor gas-flow ratio on the film properties was investigated. In particular, a wide optical band gap (2.3 eV)was reached by using a high methane-to-silane flow ratio during the deposition of the a-Si1−xCx:H layer. Theeffect of short-time annealing at 700 °C on the composition and properties of the layer was studied by X-rayphotoelectron spectroscopy and Fourier transform infrared spectroscopy. It was observed that thesilicon-to-carbon ratio in the layer remains unchanged after short-time annealing, but the reorganizationof the film due to a large dehydrogenation leads to a higher density of SiC bonds. Moreover, the film remainsamorphous after the performed short-time annealing. In a second part, it was shown that a high density(1×1012 cm−2) of silicon nanocrystals can be grown by low pressure chemical vapor deposition ona-Si0.8C0.2 surfaces at 700 °C, from silane diluted in hydrogen. The influence of growth time and silane partialpressure on nanocrystals size and density was studied. It was also found that amorphous silicon carbide sur-faces enhance silicon nanocrystal nucleation with respect to SiO2, due to the differences in surface chemicalproperties.
Silicon nanocrystals (Si-NC) embedded in insulating matrices(silicon oxide, silicon nitride) or semiconducting ones (silicon carbide)have attracted much interest for a large range of electronic devices likesingle electron transistors, silicon nanocrystal memories and morerecently “all-Si” tandem cells for third generation photovoltaics [1–3].Due to quantum confinement, the embedded nanocrystals have anadjustable band gap which can be tuned as a function of both crystalsize and host-matrix properties [3]. Therefore, it is of particular impor-tance to control and optimize the Si-NC size, spacing between Si-NC andbarrier height at the interface between nanocrystals and matrix.
The main goal of our study is to deposit amorphous silicon carbidefilms (a-Si1−xCx:H) with well-controlled properties and then to growSi-NC with controlled size and density on a silicon carbide alloy sur-face. This process can be repeated in order to obtain a multi-layerstructure containing Si-NC embedded in a silicon carbide alloy matrix.
Although Si-NC embedded in oxide [4,5] and nitride [6,7] filmshave been widely investigated, experimental investigations of amor-phous silicon carbide films with embedded Si-NC are only few [8,9].Silicon carbide alloys have the advantage of lower band gap valuescompared to the ones of silicon oxide or nitride, which means an en-hancement of the tunnelling transport between Si-NC and an increaseof the nanocomposite effective conductivity [3]. It has been shownthat the optical band gap of amorphous silicon carbide and morebroadly speaking of silicon alloys can be tailored over a wide rangeof values by controlling the composition and the nature of chemicalbonds involved in such networks [10–12]. In the aim to use Sinanocrystals embedded in a silicon carbide matrix for various opticaland electrical applications, it is of primary importance to control thematrix band gap in order to obtain a good compromise between suf-ficient quantum confinement and transport properties for the chargecarriers in the matrix.
The most widely applied technique for elaboration of Si-NC isbased on the deposition of thin films of silicon-rich material (suchas silicon-rich oxide, silicon-rich nitride or silicon-rich carbide) byplasma enhanced chemical vapor deposition (PECVD) followed byhigh temperature annealing [13–15]. In PECVD, the operating
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conditions such as RF power, total gas flow, precursor gas flow ratioor substrate temperature determine the plasma parameters, hencethe composition of the growing film [11,16,17]. The annealing athigh temperature has however some disadvantages and representsweaknesses from an industrial point of view. In particular, since therequired temperature for crystallisation of the Si nanoparticles is gen-erally above 1000 °C, only silicon wafers or quartz substrates can beused. This represents a serious limitation for low cost photovoltaicapplications.
Low pressure chemical vapor deposition (LPCVD) is a promisingway to make silicon nanocrystals due to the potential lower processtemperature. Baron et al. have shown the possibility to depositSi-NC on different dielectric substrates by controlling the early stageof the Si nucleation and growth by using LPCVD [18]. They haveobtained a high density of dots, up to 1012 cm−2 on SiO2 and Si1−xNx,and have shown that the chemical nature of the substrate has a stronginfluence on the nucleation rate.
This paper is devoted to the study of the growth by LPCVD of sili-con nanocrystals with controlled size and density on tailored siliconcarbide alloy thin films. The paper is organized as follow. The experi-mental techniques are presented in Section 2 along with the charac-terization methods of the deposited layers. Section 3 reveals theobtained results accompanied by an extended discussion on the prop-erties and the performances of the layers. Finally, Section 4 summa-rizes the main conclusions.
2. Experimental details
2.1. a-Si1−xCx:H film deposition
The a-Si1−xCx:H thin films were deposited in a single chamberPECVD system onto CZ (100) silicon substrates with thickness of525 μm, corning glass and synthetic quartz substrates of dimensions25 mm×25 mm×1 mm. Before deposition, the quartz and glass sub-strates were cleaned with soap, and then degreased in ultrasoniccleaning equipment from NOVATEC. The silicon substrates weredipped in 10% hydrofluoric acid (HF) for 10 s to remove the nativeoxide that forms on the surface.
The 13.56 MHz radiofrequency (RF) discharges were producedbetween two parallel electrodes separated by a distance of 25 mm.The gases were introduced through a showerhead upper electrode.Such a configuration ensured a good uniformity of the depositedlayer. The substrate was fixed on the bottom electrode. Mixtures ofsilane (SiH4) and methane (CH4) as precursor gases, and hydrogen(H2) were used. The methane-to-silane gas flow ratio defined asy=[CH4]/[SiH4] was varied from 1 to 14, while the total flow ratewas kept constant at 4400 sccm. During the growth process thetotal gas pressure, power density and substrate temperature werekept constant at 533.3 Pa, 162 mW/cm2 and 200 °C, respectively.
To reproduce as much as possible the temperature profile andatmosphere during the nanocrystal growth, a short-time annealingof the as-deposited a-Si1−xCx:H layer was performed in the sameLPCVD chamber as used for the Si-NC growth. The annealing isaimed at studying closely the changes in the properties of the siliconcarbide alloy induced by the process of Si-NC growth. The samplewith y=14was selected in order to analyse the main physicochemicalproperties of this kind of films before and after short-time annealingunder H2 flow for 20 s at 700 °C. Let us notice that the Si-NC weregrown on as-deposited layers.
2.2. Film characterization
Different characterization methods were used to obtain the prop-erties of silicon carbide alloy films as a function of deposition param-eters. The film thicknesses were determined by using a surfaceprofilometer from Dektat for the as-deposited layers on quartz
substrates, and by using a spectroscopic ellipsometer from Sopra forthe layers deposited on Si substrates. For all ellipsometric measure-ments, the spectra were acquired in the range from 250 nm to800 nm with a 75° angle of incidence. In order to obtain the layerthicknesses and the optical parameters, the recorded spectra of thefilms were modelled by using the Forouhi–Bloomer dispersion law.
The optical transmittance and reflectance of the films were mea-sured at room temperature by a UV–vis–NIR spectrophotometer inthe range 250–1500 nm for an 8° angle of incidence. The opticalband gap reported here was deduced from Tauc plot, which is usuallyused to describe the light absorption in amorphous and crystallinesemiconductor films.
To investigate the qualitative change in bond structures, the de-posited silicon carbide alloys were analysed by Fourier transform in-frared spectroscopy (FTIR) with a Brucker spectrometer in a spectralrange of 400–4000 cm−1 with a resolution of 2 cm−1. The elementalsilicon and carbon content of the films were obtained by energy dis-persive X-ray analysis (EDX) and X-ray photoelectron spectroscopy(XPS). XPS measurements were carried out on an Omicron MXPSspectrometer using a focused monochromatized Al Kα radiation(hν=1486.6 eV). The background pressure in the analysis chamberwas stabilized below 2×10−7 Pa during the analysis. Pass energy of10 eV enabled achieving an energetic resolution of 292 mV, in orderto optimize peak separation. No charge neutralization was used,since all recorded spectra were never shifted (with respect to the con-tamination C1s peak located at 285 eV). Appropriate ion etching wasperformed in order to eliminate the superficial oxidation and to ana-lyse the bulk material. The sputtering ion beam species are Ar+ andthe beam energy is 2 kV. The beam current density is estimated tobe 15μA/cm2.
2.3. Si-NC growth
The silicon nanocrystals growth was carried out in a cold wallLPCVD EPI Centura 5200 reactor from Applied Materials. The Si-NCwere grown onto 10 nm-thick a-Si1−xCx:H layers previously deposit-ed on silicon substrates. The Si-NC growth was performed by thermaldecomposition of silane (SiH4) in hydrogen at 700 °C under a totalgas pressure of 2.7×103 Pa. Hemispherical Si nanocrystals were as-sembled on a-Si1−xCx:H surface by using low SiH4 gas flow rate andshort growth time. Two sets of samples were prepared: for the firstset, the SiH4 gas flow rate was adjusted between 15 sccm and30 sccm. For the second set, the growth time was varied from 20 sto 35 s.
Since the PECVD reactor for silicon carbide thin film depositionand LPCVD reactor for Si-NC growth are two different tools, thewhole process cannot be done in situ. The amorphous silicon carbidefilms were treated by 10%-HF for 10 s in order to remove the surfaceoxide and to hydrogenate the surface just before the Si-NC growth.We observed that this step is essential to obtain a high density of Sinanocrystals. The HF treatment actually modifies the surface ofa-Si1−xCx:H layer through creation of large number of weak Si\H(70 kcal/mol) and C\H (97.5 kcal/mol) bonds. These bonds giverise to the following reactions at the surface during Si-NC growth byLPCVD:
SiH4ðgÞ þ Si\H→Si\SiH3ðaÞ þ H2 ð1Þand
SiH4ðgÞ þ C\H→C\SiH3ðaÞ þ H2: ð2ÞConsequently the bonding of SiH4 on the surface of a-Si1−xCx:H is
favoured and the growth of Si-NC enhanced.As the HF treatment is a non-vacuum process and according to its
importance to the high density of Si nanocrystals, an industrial solu-tion will be to use a cluster with two separated chambers dedicated
0 2 4 6 8 10 12 14 160.6
0.7
0.8
0.9
1.0
1.1
Dep
osi
tio
n r
ate
(nm
.s-1)
1 2
[SiH4] = 50sccm
y = [CH4]/[SiH4]
Fig. 1. a-Si1−xCx:H deposition rate as a function of methane-to-silane gas flow ratio.In region 1, the SiH4 flow rate decreases while in region 2 it is constant.
138 J. Barbé et al. / Thin Solid Films 522 (2012) 136–144
to the Si-NC growth by LPCVD and to the a-Si1−xCx deposition byPECVD respectively, linked by an air-lock. In this case, a multi-layerstructure could be obtained by alternating the Si-NC growth anda-Si1−xCx deposition. Such a solution will avoid the desoxydationstep since it will remain two-step vacuum process. Moreover, it willopen the possibility to hydrogenate the a-Si1−xCx surface by a hydro-gen flow just before the Si-NC growth.
2.4. Si-NC characterization
The characterization of Si-NC growth on the a-Si1−xCx:H layer wasperformed on the basis of observations by scanning electron micros-copy (SEM) and by transmission electron microscopy (TEM). Theelectron microscopy technique allows extended study in structureson nanometer scale like the Si-NC elaborated in the above describedprocess.
SEM observations were performed by using a LEO 1530 from Zeiss,equipped with a secondary electron in-lens detector. The ultimatetheoretical resolution is 1 nm at 20 kV. Before imaging in plan-view,the samples were decontaminated in plasma cleaner with a mixtureof H2 and O2 for 2 min.
Bright field TEM (BFTEM), high-resolution TEM (HRTEM) and se-lected area diffraction (SAD) experiments were carried out on a FEITechnai F30ST equipped with a field emission gun operated at300 keV (Schottky emitter). Conventional TEM sample preparationtechniques including mechanical polishing and grazing incidence(6°, 5 kV) Ar-ion milling were used to obtain samples in bothplan-view and cross sectional geometry. For TEM observation of thecross sectional samples without Si-NC, the Si substrate was orientedon a [110] zone axis. For the sample containing Si-NC, BFTEM images,HRTEM images and SAD pattern were acquired in plan-view geome-try in the vicinity of the [100] zone axis of the Si substrate. The sizedistribution of Si-NC was measured from intensity profiles onBFTEM images.
3. Results and discussion
3.1. Silicon carbide alloy
3.1.1. As-deposited a-Si1−xCx:H thin film: influence of the depositionconditions
The results presented in this section deal with the characterizationof as-deposited a-Si1−xCx:H layers. The layers are deposited in PECVDprocess by using a mixture of SiH4 and CH4 highly diluted in H2. Thedeposition conditions are summarized in Table 1. The purpose wasto select the y-ratio that gives a composition of the a-Si1−xCx:Hlayer with specific optical properties, i.e., optical band gap higherthan that of quantum confined silicon.
Fig. 1 shows variation of the deposition rate as a function ofy-values. One can clearly distinguish two regions: in region 1, for de-creasing SiH4 flow rates (see also Table 1), the deposition rate de-creases from 1 nm·s−1 to 0.7 nm·s−1. In region 2, for a constantSiH4 flow, the deposition rate is constant. Such variations indicate
Table 1a-Si1−xCx:H deposition conditions for a mixture of SiH4, CH4 and H2.
that the deposition rate is mainly controlled by the availability ofSiH4 and its radical by-products, the contribution of CH4 being lesspronounced despite the higher CH4 gas flow. Actually, Kae-Nune etal. demonstrated that SiH3 and CH3 radicals mainly contribute to thefilm growth compared to sub-radicals SiHn and CHn (nb3) [19,20].In addition, their work revealed that CH3 radical density near the sub-strate was much larger than SiH3 density for the same dissociationdegree of the parent molecule. This is due to the difference betweentheir reaction probabilities at the surface. The authors found aβ-value for the CH3 surface loss probability, which is the sum of thesticking probability and the recombination probability for a radicaldiffusing on the surface, of βCH3
=0.014 [20]. For SiH3, Matsuda etal. measured a value of the surface loss probability βSiH3
~0.26 [21].These values of β suggest that SiH3 radicals react preferentially withthe substrate and largely participate to the film-growth process,while the CH3 radicals have lower reactivity with the substrate andare more frequently desorbed into the gas phase. This explains thefact that the deposition rate is mainly controlled by the SiH4 flow rate.
The carbon content evolution relative to y=1 as determined byEDX is shown in Fig. 2a. It is difficult to prevent some carbon contam-ination at the sample surface by using EDX analysis. Therefore, thistechnique is not quite appropriate to measure the absolute carboncontent in thin films, but, at this point of the study, the relative vari-ations of the carbon content give enough information about the depo-sition process. It can be seen in Fig. 2a that the incorporated carbonfraction in the layer doubles when y ratio increases from 1 to 5, butthen tends to saturate for higher y-ratio values. This saturation ismost likely due to a lower electron density in the RF dischargewhen increasing the CH4 fraction in the mixture, as the appliedpower density and total pressure for the discharge maintenance arefixed, and the SiH4 flow is kept constant in this range of variation ofy-values. Due to the higher energies for direct and dissociative ioniza-tion of CH3 (9.8 eV and 14.3 eV, respectively [20]) compared to thoseof SiH3 (8.3 eV and 12.3 eV, respectively [20]), the density of chargedparticles in the discharge will be lower for the same applied powerdensity when increasing the CH4 component of the gas mixture.This trend of low carbon incorporation into the a-Si1−xCx:H layersis in good agreement with the observations of Pereyra et al. for similardeposition conditions [22].
Tauc band gaps of a-Si1−xCx:H layers deposited at various methane-to-silane gas flow ratios were calculated from reflection/transmissionmeasurements (Fig. 2b). The results conform well to other experimen-tal values from the literature [12,22,23] for similar conditions. In thefirst part of the curve, i.e., for y values up to 5, the optical band gap isincreased from 2.0 eV to 2.3 eV with increasing y values. For higher yvalues, the optical band gap is almost constant with a value of around2.3 eV, indicating that the film composition remains unchanged in
2.4
2.2
2.6
0
1
2
3
b
Eg
(eV
)a
C/C
0
0 2 4 6 8 10 12 14 161.8
2.0
A. Desalvo(1997)
C. Ambrosone(2002)
I. Pereyra(1996)
current paper
y=[CH4]/[SiH4]
Fig. 2. (a) Relative carbon content as a function of y measured by EDX. (b) a-Si1−xCx:Hoptical band gap as a function of methane-to-silane gas flow ratio determined by usingTauc plot. Values from other authors are depicted for comparison.
0
2
4
6
Si-Hn
Si-CH3C-Hn C-H3
Si-H
Si-O
Si-C as-deposited
annealed 700°C
500 1000 1500 2000 2500 3000
No
rmal
ized
ab
sorb
ance
Wavenumber (cm-1)
Fig. 3. Spectra of infrared absorbance normalised to film thickness obtained fromas-deposited and annealed silicon carbide alloy (SRC14).
139J. Barbé et al. / Thin Solid Films 522 (2012) 136–144
this range of y values. The results presented in Fig. 2b are consistentwith the evolution of carbon incorporation into the a-Si1−xCx:H layerdepicted in Fig. 2a: with increasing carbon content in the layer, Si\Si3Ctetrahedra are progressively replaced by Si\Si2C2 tetrahedra, thus in-creasing the number of strong Si\C bonds, which leads to an increasein optical band gap, as reported by many experimental and theoreticalworks [24–26].
These results show that the a-Si1−xCx:H deposition rate is con-trolled by varying the SiH4 flow rate while the carbon incorporation,hence the optical band gap, is controlled by varying the methane-to-silane flow ratio. For y=14, a low deposition rate and highestoptical band gap of the a-Si1−xCx:H layer were reached, which isrequired when aiming at a tailored a-Si1−xCx:H as a matrix withembedded Si-NC. The sample SRC14 mentioned in Table 1 has beenchosen for the following studies discussed in this paper, while theother a-Si1−xCx samples being with lower band gap were abandoned.Thus, we could compare the Si-NC size and density for changing pro-cess parameters on a stable and reproducible surface.
It is well known that the crystallisation of Si nanoparticles embed-ded in amorphous silicon carbide layers can be attained through hightemperature treatments [14,15]. To achieve this mechanism, tem-peratures above 700 °C are needed for the formation of siliconnanocrystals in the amorphous matrix [15]. Actually, to obtain ahigh degree of crystallisation, temperatures between 900 °C and1000 °C are necessary. In our process, the LPCVD growth of Si-NConto the silicon carbide layer surface is performed at 700 °C, but itis required to avoid crystallisation of Si nanoparticles inside thelayer since the control of Si-NC size and density would be difficultin this manner. In order to study the changes in composition, struc-tural and optical properties that can take place during the Si-NCgrowth, the silicon carbide alloys were characterized after a 20 sannealing at 700 °C.
After this short-time annealing, the layer thickness shrinks from11 nm to 7 nm. Besides, the optical band gap determined by Taucplot is reduced from 2.3 eV to 1.7 eV. The last value is rather low forapplication to solar cells. For instance, in a three-junction tandemcell, the top cell should have a band gap around 2 eV [3], whichmeans that the matrix should have a band gap above 2 eV to get aquantum confinement. However, there are no theoretical studies onthe optimized band gap for the matrix, but it might be above 2 eVto get a quantum effect and below 2.5 eV to assure a good conductiv-ity. In this study, we show that we can fine-tune the band gap of thematrix. In a future work, the matrix properties will be optimized andparticular attention will be paid to the band gap in order to select themost appropriate one for application to solar cells. From the followingFTIR and XPS characterizations, some explanations based on chemicalstructure modifications will be given for the above mentioned effects.
In order to understand the qualitative change in bond structures,the infrared absorption spectra before and after annealing normalisedto the respective film thicknesses, i.e. considering the film thicknessshrinkage during annealing, are given in Fig. 3. The broad peak locat-ed at 2100 cm−1 with a Gaussian line shape, which is related toSi\Hn stretching bonds is absent after annealing. This indicates alarge out-diffusion of hydrogen. We observe the same effect for theweaker peaks related to Si\Hn wagging modes at 640 cm−1 and670 cm−1, CH3 stretching modes at 2880 cm−1 and 2960 cm−1,CHn wagging mode at 1000 cm−1 and Si\CH3 bending mode at1240 cm−1. Different authors showed the dissociation of hydroge-nated bonds (Si\H, C\H) for high temperature treatments [27–29].Due to hydrogen out-diffusion, it has also been observed that the re-sidual Si and C atoms rearrange to form new Si\C bonds which re-sults in the increase of the Si\C peak intensity [27]. It can be seenin Fig. 3 that the Si\C peak intensity is increased by a factor 3 afterannealing. Let us note the broad full width at half maximum(FWHM) induced by annealing, which is probably a signature oflarge disorder in the SiC network (Si\C length and Si\C\Si angle).The peak at 1107 cm−1 is attributed to Si\O bonds in a carbide envi-ronment [30].
It is generally acknowledged that after crystallisation of SiC, theshape of Si\C absorption peak changes from Gaussian type to Lorentztype. Moreover, the peak is blue-shifted and the FWHM is reduced asan indication of the transition from amorphous to crystalline state ofthe film [31]. None of these observations has been noticed in our film.This indicates that SiC nanocrystallites do not form at 700 °C. Howev-er, the infrared spectra do not give any information about thecrystallisation of excess silicon.
The large decrease in hydrogen content can explain the layershrinkage and the decrease in the optical band gap after heat treat-ment. Indeed, addition of hydrogen decreases the density of localised
Co
un
ts (
a.u
.)
as-deposited
700°C annealed
CHn
CSi
Si-CHnC1s
292 290 288 286 284 282 280Binding energy (eV)
Fig. 5. Narrow-scan XPS spectra of C1s core level acquired after ion etching fromPECVD silicon carbide alloy (SRC14) as-deposited and annealed at 700 °C for 20 s.
140 J. Barbé et al. / Thin Solid Films 522 (2012) 136–144
states and thus increases the optical band gap [32]. It seems that thecompetition between the increase in Si\C bond density, whichshould increase the optical band gap, and the hydrogen out-diffusion is in favour of the later one.
Values of carbon and silicon concentrations were estimated fromXPS. Fig. 4 shows wide-scan XPS spectra of as-deposited and annealedsilicon carbide alloy. In order to get rid of carbon and oxide surfacecontamination of layer due to air exposure, Ar+ etching wasperformed, up to 10 min at 2 kV. This also allowed reaching the filmbulk until we got a constant Si/C level. The Ar2s and Ar2p peaks onthe wide-scan spectra are due to a slight implantation of Ar atomsduring ion etching. The as-deposited film composition estimatedfrom the peak intensity is 69.6 at.% silicon and 21 at.% carbon, withsome oxygen contamination (9.4 at.%). Since oxygen atoms arefound in the bulk of the layer, and not only at the surface, this con-tamination has most probably taken place during the thin film depo-sition process itself. After annealing, the composition is almostunchanged with 70.3 at.% silicon, 18.7 at.% carbon and 11 at.% oxy-gen. These data show that the deposited layer is therefore asilicon-rich carbide material and that 700 °C short-time annealinghas no effect on the Si, C and O composition.
C1s high resolution scans before and after annealing are repre-sented in Fig. 5. Before annealing, the main carbon peak located at283.8 eV reveals the presence of carbon in an intermediate com-pound (Si\CHn) between pure SiC structure (CSi at 283 eV), andamorphous hydrogenated carbon (CHn, from 284 to 286 eV). Afterannealing, the main carbon peak is shifted to 283 eV (CSi). This ener-gy corresponds to SiC with a C atom bonded exclusively to Si atoms.This is in good agreement with the FTIR observations: hydrogenout-diffuses during annealing leading to a restructuration of thelayer. It means that Si and C dangling bonds are formed by the H de-hydrogenation. Owing to the low concentration of carbon, C danglingbonds react with Si dangling bonds only to create new SiC bonds in anamorphous Si matrix. It is clearly revealed by the peak at 283 eV.
From these results it can be assumed that, after 700 °C short-timeannealing, the film is homogeneously composed of 37.4% SiC mixed ina Si matrix with a low SiO content. This conclusion is in line with theresults for the optical band gap and gives explanation of the quitehigh value of the optical band gap for high y-values. For the sake ofclarity, the thin film will be defined as a-Si0.8C0.2 in the following ofthis paper.
Cross-section HRTEM observations were performed to reveal themicrostructure of the annealed film and the interface morphology be-tween the film and the substrate. Fig. 6 shows the a-Si0.8C0.2 film onsilicon substrate after 700 °C short-time annealing. First, a clear dis-tinction in structure is observed between the substrate and the film:after annealing, the Si substrate shows crystalline planes while thea-Si0.8C0.2 film remains amorphous. SAD pattern (shown in the inset
Si2p
Si2s
Ar2p
C1s
Ar2s
Co
un
ts (
a.u
.) OKLL
O1sas-deposited
annealed 700°C
1000 800 600 400 200 0Binding energy (eV)
Fig. 4.Wide-scan XPS spectra acquired from as-deposited and annealed silicon carbidealloy (SRC14) after ion etching.
of Fig. 6) acquired from the same processed materials, but inplan-view geometry, confirms that the a-Si0.8C0.2 film remains amor-phous after short-time annealing: no silicon nanocrystals wereformed inside the film which would be evidenced by diffractionrings. These results are not in accordance with Künle et al. who ob-served that excess of Si starts to crystallise in a-Si0.8C0.2 layers at tem-peratures of about 700 °C during annealing under N2 for 30 min [15].However, this discrepancy with our results could be explained by dif-ferent annealing conditions. In our case, we use a rapid thermalannealing under H2 for only 20 s. This is most likely the main reasonfor the obtained results. Second, we notice that the interface betweenthe film and the polished wafer has a roughness in the nanometerrange, which may indicate a slight epitaxial crystallisation of siliconexcess in the film nearby the substrate. The SAD pattern in Fig. 6indicates that the excess silicon is grown along the crystalline orien-tation of Si substrate. Nevertheless, if the latter can promote thecrystallisation of silicon excess near the interface, it has no influenceon the bulk of the a-Si0.8C0.2 film which remains amorphous afterour 700 °C short-time annealing. Another hypothesis which could ex-plain the rough interface is hydrogen plasma etching of the siliconsubstrate during the a-Si0.8C0.2:H deposition. This is most likely dueto the high H2 dilution rate and relatively high power density(162 mW/cm2) used during the deposition. Furthermore, we haveobserved that the deposition rate is twice higher if argon is used in-stead of hydrogen as diluting gas (results not shown in the currentpaper). It might indicate that hydrogen etches silicon during the de-position (the substrate as well as the thin film). So, epitaxialcrystallisation and etching are two mechanisms acting complementa-ry and rather correctly explaining the interface roughness.
In the first part of this paper, the deposition conditions for theelaboration of a-Si1−xCx:H films were studied. A silicon-rich siliconcarbide layer was characterized before and after a short-time (20 s)annealing at 700 °C in order to study the changes that can takeplace during the Si-NC growth by LPCVD. It was found that theannealing leads to H out-diffusion associated with a decrease of theoptical band gap and formation of Si\C bonds. However, theannealing does not provoke the formation of Si-NC inside the siliconcarbide thin film which remains amorphous.
3.2. Silicon nanocrystals
3.2.1. Si-NC on a-Si0.8C0.2The second part of this paper deals with the growth of Si-NC by
LPCVD on as-deposited 10 nm-thick a-Si0.8C0.2 thin films. SampleSRC14 was still used for this study, so that the growth takes placeon a well-defined and controlled layer. The Si-NC growth conditionsare summarized in Table 2, along with the Si-NC surface density andsize. The Si-NC surface density was obtained from the SEM images
Fig. 6. HRTEM cross-section image and SAD pattern from plan-view samples (inset) of a-Si0.8C0.2 layer on Si substrate after 700 °C short-time annealing.
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by direct counting. The size of nanostructures was evaluated from thesame images. This method admits an error of 5% on the density valueand from 15% to 50% on the dot size (the smallest dots being the mostdifficult to evaluate). Fig. 7 presents a plan-view SEM image of sampleQDf30 for which a high Si-NC density of 1.1×1012 cm−2 wasmeasured.
In the aim to control the Si-NC density and size, two sets of sam-ples with various SiH4 flow rates and growth time were prepared,as described in Table 2. It can be observed in Fig. 8 that the Si-NC den-sity increases as the growth time (Fig. 8a) or the SiH4 flow rate(Fig. 8b) increases. We observe the same variations for the Si-NCmean diameter: all along the process, the Si-NC grow until coales-cence for either high SiH4 flow rate or long growth time. The highestachievable Si-NC density was 1.1×1012 cm−2 for a mean size of6 nm. For low SiH4 flow rate or short growth time, the density andsize of Si-NC were difficult to estimate due to the limited SEM resolu-tion. These trends are similar to those described by Mazen et al. forgrowth of Si-NC on SiO2 films [33]. Nevertheless, the growth timeand SiH4 flow rate ranges are much lower in our case. Furthermore,high Si-NC densities were obtained even for very low values of pro-cess parameters (samples QDt20 and QDf15).
3.2.2. Influence of the surfaceIn order to study the effect of surface chemical properties on the
Si-NC nucleation, Si-NC were grown on a-Si0.8C0.2 and on SiO2 layersby using the above described growth technique in the same processconditions for both surfaces. SiO2 has already been studied and highSi-NC densities up to 1012 cm−2 were reached on highly hydroxylat-ed SiO2 surfaces [34]. For SiO2 surfaces, the study of nucleation ratehas shown that hydroxylated SiO2 surfaces favour the nucleation ofsilicon with respect to non-hydroxylated SiO2 due to the lower ener-gy of OH bond compared to the Si\O bond [34]. Here, 300 nm-thickSiO2 film was thermally grown on (100) Si substrate at 950 °C. HFsurface cleaning was performed on both a-Si0.8C0.2 and SiO2 samples,with 10% HF for 10 s just before Si-NC growth. The growth time was15 s for both samples. The variation of Si-NC density grown on
Table 2Evolution of the nanocrystals density and size for the two sets of sample.
a-Si0.8C0.2 and SiO2 surfaces as a function of SiH4 flow rate is repre-sented in Fig. 9.
One can observe in Fig. 9 that, for the same growth time and SiH4
flow rate, the Si-NC density is almost 10 times higher for Si-NCgrowth on a-Si0.8C0.2 layer than on SiO2 layer. Three distinct reasonsare responsible for the higher Si-NC nucleation rate on a-Si0.8C0.2compared to the SiO2 surface. First, Si\H and C\H bonds have weak-er bond energies (70 kcal/mol and 97.9 kcal/mol, respectively) thanthe O\H bond energy (102 kcal/mol). It implies that Si\H andC\H bonds are more easily broken via chemical reaction withadsorbed silane radicals than the O\H bond, particularly if the sur-face is silicon rich, like for the present a-Si0.8C0.2 layer. In the case ofa-Si0.8C0.2, the HF surface cleaning performed before each depositionleads to desoxydation and hydrogenation of the surface. Thehydrofluoric acid etches completely the siloxane groups (Si\O\Si)and leads to the formation of Si\H and C\H bonds. In the case ofSiO2, the HF surface cleaning is meant to break Si\O bonds, to partlyhydrogenate the surface and to create silanol groups (Si\OH). Thus,the SiO2 film is partly terminated by silanol groups while thea-Si0.8C0.2 film has a majority of Si\H and C\H terminations on thesurface that promote the Si-NC nucleation by chemical reactions.
Second, it has been observed that PECVD deposited thin films en-hance the Si-NC nucleation compared to thermally grown SiO2 [35],which is due to a higher degree of dangling bonds and defects atthe layer surface, as well as to a higher surface roughness.
Third, the a-Si0.8C0.2 layer has a higher silicon concentration (90%mass concentration) than the silicon oxide layer (46% mass concen-tration). The silicon atoms are known to be preferential nucleationsites for silane radicals [36], which makes the silicon rich carbidelayer a preferential surface for the nucleation of silicon nanocrystals.
Fig. 7.Plan-view SEMobservations of Si nanocrystals grownon a-Si0.8C0.2 (sample QDf30).
Fig. 8. Variations of density and mean size of Si-NC grown on a-Si0.8C0.2 at 700 °C as afunction of growth time (a) and SiH4 flow rate (b). Lines are shown as guide to the eye.
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As far as the growth mechanism is concerned, one can considerthe Si-NC nucleation mode on a-Si0.8C0.2 to be half-way in-betweenthe Frank–van der Merwe growth and Volmer–Weber growth, butdifferent from Stranski–Krastanov growth [37]. The adatoms do notdiffuse much on the surface, but attach favourably both to the surfaceand to the silicon clusters. Indeed, the a-Si0.8C0.2 layer is silicon richand so behaves closely to silicon clusters. On the other hand, it hasbeen shown that the nucleation of silicon on SiO2 surface is the
1
10
a-Si0.8C0.2
SiO2
Den
sity
(10
11 cm
-2)
15 20 25 300.1
SiH4 flow rate (sccm)
Fig. 9. Variation of Si-NC density on thermally grown SiO2 and PECVD depositeda-Si0.8C0.2 as a function of silane flow rate. Growth time is 15 s in both cases. Linesare shown as guide to the eye.
Volmer–Weber type of nanocrystal growth [33]. In this second case,the adatoms diffuse a lot on the surface to finally attach to otheradatoms or silicon clusters, which explains the lower nanocrystalsdensity at the same process conditions.
3.2.3. TEM observationsIn order to better understand the size distribution and the crystal-
line structure of Si nanoparticles, TEM based techniques, such asHRTEM, BFTEM imaging and SAD were applied in plan-view samplegeometry (sample QDf30). TEM is the only available techniquewhich gives direct information about the crystalline structure of theSi-NP and accurate determination of their size. In Fig. 10a, Sinanoparticles appear as dark circular contrast in the BFTEM imagewith a density of 1.3×1012 cm−2. The density found from TEM mea-surements is slightly higher than the nanoparticle density evaluatedfrom SEM images which was found to be about 1.1×1012 cm−2.This is consistent since closely spaced nanoparticles will be countedas one particle due to the lower resolution of the SEM. From thesame image, we determined the size distribution for a sample groupof 49 nanoparticles. Only isolated nanoparticles were taken into ac-count in this study although a significant amount of agglomeratesare visible on the surface. The histogram in Fig. 10b gives a mean di-ameter of 4.8 nm and a standard deviation of the nanoparticle size of0.7 nm. These values are consistent with the value of 6.2 nm foundfrom SEM images, given the uncertainty of the SEM measurements.Fig. 10c shows high resolution HRTEM plan-view image of one nano-particle at high magnifications. The nanoparticle has (200) latticeplanes visible. In the SAD pattern shown in Fig. 10d, Si-NC reflectionsof Si (200) and (111) lattice planes are observed. It indicates thatSi-NC do not have any preferential orientation with respect to thesubstrate plane. These observations clearly show that the siliconnanoparticles grown on a-Si0.8C0.2 thin films by LPCVD at 700 °C arecrystalline.
4. Conclusion
a-Si1−xCx:H thin films were deposited by PECVD in SiH4, CH4 andH2 mixtures. The carbon content, optical band gap and depositionrates were tailored by testing various methane-to-silane flow rates.Optical band gap of 2.3 eV and carbon content up to 20% werereached in the as deposited films. Short-time annealing of the filmat 700 °C leads to the out-diffusion of hydrogen, as well as to thinfilm shrinkage and to optical band gap decrease. The Si and C compo-sitions given by XPS are found to remain unchanged after short-timeannealing but the film rearranged during annealing process, thus in-creasing the Si\C bonds density. However, the annealing does notlead to the formation of silicon nanocrystals inside the silicon-richsilicon carbide alloy, which remains amorphous.
Silicon nanocrystals were subsequently grown by LPCVD at 700 °Cfrom SiH4 diluted in H2 on as-deposited a-Si0.8C0.2 thin films. Weobtained Si-NC densities higher than 1×1012 cm−2, even for lowgrowth times and low SiH4 flow rates. Compared to SiO2 surfaces,a-Si0.8C0.2 layer is much more appropriate to obtain higher Si-NC den-sity for same process conditions. This affinity is explained by surfacechemical composition, structural quality and high silicon concentra-tion at the surface. The obtained results reveal that amorphous siliconcarbide films are promising materials for the growth of siliconnanocrystals on a semiconducting matrix with tailored properties.
Acknowledgment
The authors acknowledge the support from the EU funded FP7Project “SNAPSUN”.
Fig. 10. Plan-view TEM observations of Si-NC grown on a-Si0.8C0.2 surface (sample QDf30): (a) BFTEM image, (b) size distribution determined from panel a for isolated Si-NC,(c) plan-view HRTEM image at high magnification. (200) and (111) planes of Si nanocrystals are observed, and (d) SAD pattern shows the diffraction pattern of Si substrate on[100] zone axis as well as Si (200) and Si (111) reflections from the silicon nanoparticles.
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