E L S E V I E R Thin Solid Films 295 (1997) 37-42

Effect of chamber pressure on p-type/xc-SiC:H thin films prepared by photo-CVD

Arup Dasgupta a, Sukriti Ghosh a, S.T. Kshirsagar b, Swati Ray a,. a Energy Research Umt, Indian Assoctationfor the Cultivation of Science, Jadavpur, Calcutta 700 032, India

b Physical Chemistry Division, Nattonat Chemtcal Laboratory, Pune 411008, India

Received 15 February 1996; accepted 22 August 1996

Abstract

Highly conducting boron-doped microcrystalline silicon carbide (#c-SiC:H) thin films have been prepared by mercury sensitised photo- chemical vapor deposition. The chamber pressure was identified as one of the most crucial parameters governing the microcrystalline growth as well as the dopant incorporation in the microcrystalline thin films. Raman studies show the crystalline size and volume fraction decreases with increasing pressure. Transmission electron microscopy of such films reveal that the crystalline phase contains silicon only, so that carbon is incorporated only in the amorphous phase. The presence of carbon in these films was confirmed by secondary ion mass spectroscopy. There is an optimum pressure (0.5 Torr), depending upon the other deposition parameters, for which the conductivity of p-type #c-SiC:H film is highest (1.2 × 10 .4 S cm- i). © 1997 Elsevier Science S.A.

Keywords" Chermcal vapour deposltmn; Silicon carbide; Electrical properties and measurements; Structural properties

1. Introduct ion

Hydrogenated silicon carbide, both amorphous and micro- crystalline are important materials for photovoltaic devices. The wide gap and good valency electron controllability [ I ] are two of its most attractive properties. Boron-doped amor- phous silicon carbon alloys (a-SixC:_x:H) are used as the window layer in amorphous silicon p - i -n solar cells [2]. Carbon incorporation increases the optical gap of p-type hydrogenated amorphous silicon (a-Si:H) thin films (its optical gap can be higher than 2 eV), and consequently lowers the conductivity by several orders of magnitude [3 ]. Micro- crystalline silicon carbide films (/xc-SiC:H), on the other hand, simultaneously offer higher electrical conductivity and bandgap than their amorphous counterparts and, therefore, is the best choice amongst the wide bandgap materials for window layers of a-Si:H based solar cells.

The plasma-enhanced chemical vapour deposition (PECVD) of/xc-SiC:H films grown from Sill4 + CH4 and Sill , + CH3SiH3 source gas mixtures has been reported [4- 7 ]. These films are composed of submicron size Si crystallites embedded in an a-SiC:H matrix. High r.f power density ( > 100 mW cm -a) together with high hydrogen dilution of source gases are the general prerequisites for growing/xc-

* Corresponding author

0040-6090/97/$17.00 © 1997 Elsevier Science S.A. All nghts reserved P I I S 0 0 4 0 - 6 0 9 0 ( 9 6 ) 0 9 2 8 2 - 6

SiC:H films [4-6]. High hydrogen dilution assists the micro- crystallite formation through high surface coverage or etching action [8-11]. But high r.f (radio frequency) power gener- ates highly energetic ions and electrons. Consequently, tex- tured transparent conducting oxide (TCO) substrates, onto which the window layer of solar cells are grown, may not be able to withstand the impact of such energetic charged par- ticles [ 12] and are thus susceptible to damage. Hence there is a need to search for an alternative technique, where energetic particle bombardment could be mmimised or eliminated.

We have recently reported the preparation of p-type/zc- SiC:H [ 11 ] using high hydrogen dilution but at a much lower power density ( ~ 10 mW cm -2) in a mercury sensitized photo-chemical vapour deposition (photo-CVD) system. This system is free from particle bombardment and, under optimised conditions, conductivity as high as 1× 10 .2 S c m - : for such films has been achieved [ 13]. This suggests that the use of photo-CVD alone would be beneficial.

In the photo-CVD technique, ultraviolet light ( 185 nm and 254 nm) is used to excite and dissociate the reactant gases. The photon energy (6.71 eV and 4.89 eV) is expected to have an effect on the surface nucleation and hence assist crystallisation since the energies are greater than the work function of silicon (4.1 eV) [ 14]. As only atomisation and

38 A Dasgupta et al /Thin Sohd Films 295 (1997) 37-42

no ionisation of the gas-phase reactants is involved in this technique, the films thus grown are not beset with ion damage.

A crucial parameter in the mercury sensitized photo-CVD system, is the pressure of source gases in the chamber [15,16]. Moreover, the acetylene gas (C2H2) has a high photodissociation cross-section producing C2H at low pres- sure [ t7]. The influence of these sensitive parameters on the growth of p-type /xc-SiC:H from a gas mixture of Sill4 + C2H2 + BaH6 + H2, hence forms the subject of the present investigation. Attempts were therefore made to exam- ine and correlate the electrical and optical properties with the structural characteristics of the films prepared at differ- ent chamber pressures; moreover, the film properties are discussed from the standpoint of growth kinetics.

2. Experimental

The details of deposition of /xc-SiC:H films using the SAMCO (Japan) mercury sensitized photo-CVD system are described elsewhere [ 11 ]. Silane ($1H4). acetylene (C2H2), diborane (BAH6) (diluted in hydrogen (1%) ) and hydrogen (I-I2) were used as source gases with independent mass flow rates of 5 sccm, 0.2 sccm, 0.7 sccm and 150 sccm, respec- tively. H e was used to carry mercury (Hg) vapour into the reaction chamber by passing it over a Hg pot maintained at 70 °C. The desired chamber pressure (PT) was maintained constant by controlling a throttle valve manually and was measured by a capacitance manometer (MKS Inst.. USA). PT was varied from 0.3 Torr to 0.8 Torr ( 1 Torr = 133 Pa) while the substrate temperature was always kept at 250 °C.

The substrates used for film deposition were Coming 7059 glass for optoelectronic and Raman studies; carbon-coated copper grids for transmission electron microscopy; and c-Si for secondary ion mass spectroscopy (SIMS). The film thick- nesses were measured by a stylus type instrument (Planar products, UK). Dark and photoconductivities (under 50 mW cm-2 white light illumination) were measured after anneal- ing the samples in vacuum ( ~ 10- 6 Torr) at 150 °C for about an hour which was expected to remove adsorbed moisture and hght induced defects, if any. A Hitachi 330 UV-VIS spectrophotometer measured the optical absorption spectra.

3. Background information

In a mercury sensitised photo-CVD system, the heated mercury (Hg) vapour is carried into the reaction chamber by hydrogen gas. These Hg atoms are then spectroscopically raised to their excited state (Hg*) by LrV photons (185 nm and 254 nm), and subsequently they collide with gas mole- cules to produce the reactant species. The film deposition then takes place through a series of primary and secondary reactions involving silane, acetylene, diborane and hydrogen gas mixed together with Hg*,

H g + h v ~ H g *

Hg* + Sill4 ~ Hg + Sill 3 q- H

Sill 3 -t- Sill 3 --~ Si2H 6 or Sill4 + Sill 2

C2I-I2 + hv ~ C2H + H

Hg* + Call2 ~ Hg + Call + H

Hg* + BaH6 ~ B 2I-I5 + H + Hg

B 2 H s + H ~ 2 B H z + (3 -Z)H2 z - - -0 -3

Hg* + H z ~ H g + 2 H

SiI-L + H ~ Sill3 + H2

Sill3 + H ~ Sill4

(1)

(2a)

(2b)

(3)

(4)

(Sa)

(5b)

(6a)

(6b)

(6c)

In all the above reactions, the collision probability of Hg* with the gas molecules plays an important role. This proba- bility depends strongly on the chamber pressure (PT) which determines the mean free path of the gas molecules. PT thus becomes a crucial parameter for the mercury sensitisedphoto- CVD system. Besides, in reaction (3), although no collision with Hg* is involved as such, the photo-dissociation cross- section of acetylene is higher for lower PT ( < 1 Torr) [ 17]. The direct photolysis of Sill4 has been reported to be negli- gible [ 18,19]. In fact, direct UV excitation of Sill 4 requires high energy photons of h u > 8 eV (h < 155 nm) [ 19 ]. More- over, preparation of p-type SiC:H films by direct photo-CVD was reported earlier [20] which proves that diborane can be dissociated by UV alone yielding BH~ (z = 0-3) type of rad- icals as shown in reaction 5(b). However at low pressures in mercury sensitised photo-CVD, reaction 5 (a) is vital [ 11 ].

4. Results and discussion

Fig. 1 shows the variations of darkconductivity (O'o) ,pho- toconductivity (Crph) and activation energy (AE) of dark

1¢ ~ 1.o

/ / ' 0.9

lOS ' 0.8 16 6

"' / 0.7

b 16 8

16,o ', I o.4

i(~I ~ / ' " " , , ~ / /

o2 o'3 o'~ o'~ o16 o'7 o0 oI~ 2 PT (Tort)

Fig I. Variations of conductlvlties (dark (o'd) and photo (o-ph), ); and acuvatxon energy of dark conducttvity( A E, - - -) with chamber pressure (Pvm Torr; 1 Torr= 133 Pa) of/xc-SxC:H thin films,

A Dasgupta et al . /Thm Solid Films 295 (1997) 37-42 39

conductivity with changes in PT. As PT is increased from 0.3 Torr to 0.5 Torr, the dark conductivity increases from 8.4× 10 .9 S cm -1 to 1.2× 10 .4 S cm -1 i.e. an increase by five orders of magnitude. As PT is increased further to 0.8 Torr, there is a sharp fall in cr d to ~3.85 × 10 -1°" S cm -1, 1.e. the decrease is nine orders of magnitude. O'ph follows a similar trend as o'a but with a much smaller gradient of increase or decrease. The AE curve is almost a mirror image of the cra curve.

The low Crph value of the film prepared at PT=0.30 Torr (O'ph/crd~-10 only) suggests that the sample is possibly microcrystalline although o'a is fairly low. Moreover, the high AE value of 0.57 eV indicates that the sample was inade- quately doped. This may be attributed to a combination of two reasons: (1) low boron content (which necessitates ele- mental analysis for such films), and (2) a lower amount of BHz radical formation wa secondary reaction (5b) because of the higher mean free path of B2H5 radicals at low pressure. The B2H5 radical can thus reach the film growing surface directly, which may not lead to efficient doping. With increas- ing PT, era reaches a maximum of 1.2 × 10-4 S c m - ~ at 0.50 Torr, where AE exhibits a minimum value of 0.25 eV. The high conductivity suggests that the film was not only micro- crystalline, but also was doped efficiently (AE is a mini- mum), because of an abundance of BH~ (z=0--3) type of radicals. The film was not photoconductive at all, as is com- mon to the family of doped microcrystalline films.

The drastic decrease in crd for PT > 0.5 Torr coupled with an increase of photoconductivity gain (O'ph/O'd) points towards a possible structural change from microcrystalline to amorphous, or in other words, a reduction in the fraction of microcrystalline phase and increase of amorphous fraction. Such changes in conductivity have been observed by other workers during structural change in the film [ 10,21 ]. As the total gas flow rate remains unchanged, an increase of PT results in an increase of precursor density and an increase in intermolecular collision probability which in turn results in increase of the precursor residence time. This may initiate some secondary reactions in the gas phase, and higher order hydrocarbons may form introducing disorder in the film, in turn destroying the crystallinity.

Saxena et al. [ 22] in their gas-phase modelling on a photo- CVD reactor, state both high hydrogen ddution and low pres- sure are necessary to ensure a high atomic hydrogen flux near the film growth surface. This leads to nucleation and conse- quently growth of microcrystallites through an etching mech- anism. They proposed that, at a constant hydrogen dilution ratio, the atomic hydrogen flux at the substrate at first increases with pressure but beyond a certain pressure (opti- mum PT) falls off quite fast. In fact, beyond the optimum pressure point, low diffusivities cause almost all hydrogen atoms, formed by reaction (6a), to react in the gas phase with Sill4 (reaction (6b)) or even $1H3 radicals (reaction (6c)) before diffusing to the growth site, thereby depleting the growth site from the key factor instrumental in microcrystal- lisation [ 11]. Therefore, a structural change is expected

2.7 220

2.6 / " / ~ - - 1so

25

r,r 100 /

23 ,fl, /"

/ " ~ 6.0 22 ..c~" o

0..-~

21 2 0'3 0'4 0~5 0'6 0~7 0'8 0920 Pr[Torr)

Fig. 2. VariationsofE04 (energy value at which c~= 1 × 104cm -I) ( ), and growth ra te (Rg) ( - - - ) , with chamber pressure (PT in Torr, 1 Torr= 133 Pa)

beyond this optimum reactor pressure point. In the present case this value corresponds to 0.50 Torr. As PT ]s increased to 0.8 Torr, the photoconductivity gain is more than three orders of magnitude in&caring that the film is amorphous. Consequently the A E value is as high as 0.95 eV (easy doping sites, i.e. crystallites are few or absent).

Fig. 2 describes the variation of Eo4 [23] (a measure of the optical gap of microcrystalline films, since Tauc's gap is not valid for them) and film growth rate (Rg) with changes in reactor pressure. Eo4, derived basically from the optical absorption spectra, is the energy value at which a ~ 1 × 10 4

cm- I . Thus No4 reflects two distinct parallel segments (one between 0.3--0.5 Torr and the other between 0.6-0.8 Torr) along with a sharp change of slope while going from 0.5 to 0.6 Torr. This is suggestive of two different classes of sam- ples. Thus the sharp change of slope can be designated as the signature of a structural change which must have occurred somewhere between 0.5 and 0.6 Torr of PT. However, no intermediate pressure point could be studied in the 0.5--0.6 Torr range, as there was a limitation in the controllability of the throttle valve. Such a structural change is in agreement with the drop in the cr d value seen in Fig. 1 where we have conjectured a change of phase from microcrystalline to amor- phous to occur when PT was increased beyond 0.5 Torr. Variation of Rg on the other hand, is almost a mirror image of the Eo4 curve. The film growth process may be an outcome of a competition between deposition of Si containing species into microcrystalline and amorphous surface sites and etching of the amorphous phase [22] i.e., deposition rate (amorph. + cryst.) - etch rate (amorph.) = growth rate. At low pressures (PT < 0.5 Torr), the film growth process is etch-rate limited. As already explained, at such pressures, the film etchants, namely atomic hydrogen flux, together with the highly reactive C2H radicals covering the film growth site is high and hence Rg is low. But at pressures beyond the optimum (PT > 0.5 Torr), not only the probability of for- mation of C2H radicals is low [ 17] but also the hydrogen

40 A Dasgupta et aL ~Thin Solid FErns 295 (1997) 37-42

flux covering the film growth surface decreases rapidly [ 22]. In other words, the growth rate is no longer etch-rate limited. This together with the presence of a higher precursor density in the reactor brings about a sudden increase in the growth rate and a subsequent destruction of the mlcrocrystallinity.

Next, structural information about the films was sought from the Raman spectra and transmission electron microg- raphy (TEM) and diffraction (TED) patterns of the films prepared at the lowest pressure of 0.3 Torr and optimum pressure of 0.5 Torr. Fig. 3(a) and 3(b) shows the Raman spectra, in the 450-550 cm- 1 range, of the samples prepared at 0.3 Torr and 0.5 Torr, respectively. The spectra show a couple of peaks identified as transverse optic (TO) hke modes of c-Si and a-Si [24]. A computer program was used to fit and deconvolute these two peaks. For both the samples, the a-Si peak was broadly distributed around 490 cm - 1 com- pared to 480 cm - 1 for very good a-Si:H films. Such Gaussian broadening around 490 cm -1 instead of 480 cm-1 was reported earlier by Kshirsagar et al. [25]. The c-Si peak normally occurs at 520 cm - 1 [ 24 ]. However, in our samples, Fig. 3(a) shows the peak center was at 518 cm -1 (PT=0.3 Torr) while in Fig. 3 (b) it was at 517 cm- ~ (PT----- 0.5 Torr). Such deviations from 520 cm- ~ owe their origin to varying crystallite size [26,27]. From the peak position of the crys- talline part in the Raman spectra, an estimate of the particle size was made from the studies of particle size variations m the range of 512-520 cm-1 by Iqbal and Veprek [26]. The particle size for 0.3 Torr deposited sample was estimated to be between 80 and 100 ,~ while that for 0.5 Torr deposited sample was smaller, between 60 and 70 It. The volume frac-

35 (a)

30

25

5

0 " - " 450 470 490 510 530 550

Raman Shlf'l( crn "I ) 12 5 (b)

IO.O

i

2.5

0 450 470 490 510 1 B30 550

Roman Shift(era- )

Fig 3. Raman spectra of/zc-SiC:H than films prepared at (a) PT = 0 3 Torr and (b) P-r=0 5 Torr ( 1 Torr= 133 Pa). The transverse optic (TO) vibra- tional peaks corresponding to c-Sx and a-Si are deconvoluted and shown by dotted lines

tion of crystallites for both the samples was estimated using the relation Pc = It(It +y I=); where Ic and I~ represent the integrated Raman intensity corresponding to crystalline and amorphous phases, while y (taken as 0.7 here) is the ratio of the integrated Raman cross-section of c-Si to a-Si [28,29]. Pc was estimated to be nearly 60% and 54% when the chamber pressures were 0.3 and 0.5 Tort, respectively. We, therefore, conclude that, as long as microcrystallinity is sustained, an increase in pressure has a inverse effect on crystallite size as well as the volume fraction. In other words, the amorphous phase fraction increases with increase of PT. Furthermore, the reason behind the poor doping of the film prepared at 0.3 Torr (A E= 0.57 eV) compared to that prepared at 0.5 Torr (AE-- 0.25 eV) in spite of bigger crystallite size and higher crystalline volume fraction for the former, was surely due to lower concentrations of efficient dopant species present in the chamber at this pressure, as inferred from Fig. 1.

Fig. 4 shows the TEMs and corresponding TEDs of the two microcrystalline samples discussed above together with the TED of the sample prepared at 0.6 Torr. The well-defined concentric diffraction rings (Fig. 4(b) and4(d)) correspond to crystalline planes, namely (111), (220) and (311) ofc- Si. Apparently, carbon was incorporated only in the grain boundaries and amorphous parts of the films. Thus the/xc- SiC:H samples were composed of Si crystallites embedded in an a-SiC:H matrix as reported by most researchers in this field [ 11,16,30,31 ]. The particle size as observed from the TEMs in Fig. 4(a) and 4(c), are much larger (500-900 in Fig. 4(a) and 100-800 * in Fig. 4(c) ) than those esti- mated from Raman spectra. Thus each microcrystal may actu- ally be a conglomeration or clustering of a number of small single crystallites. Such a phenomenon is common to micro- crystalline thin films [28]. A more careful observation of the TEMs in Fig. 4(a) and 4(c), reveals the grains are more closely and uniformly distributed in the film prepared at 0.5 Torr in which one can visibly notice the grain boundaries. As the resistivity of the inter-grain amorphous matrix is much higher than that of the crystallites, one expects a higher mobil- ity of the charge carriers in the film prepared at 0.5 Torr than at lower pressures. However the TED in Fig. 4(e) shows diffused halo, characteristic of amorphous materials. This supports the structural change during change of PT from 0.5 to 0.6 Torr and explains the conductivity variation shown in Fig. 1.

Elemental analysis for Si, C, B and H in boron-doped silicon carbide films prepared at different pressures were probed by secondary ion mass spectroscopy (SIMS) and the data are presented in Table 1. The table shows that the boron incorporation in the films are lower at lower pressures. But the SIMS result does not exactly give the percentage of boron atoms that actually goes into four-fold coordination with sil- icon. However, it will not be unjustified to assume that boron doping of silicon will be poor where boron incorporation itself in the film is low, so long as there occurs no structural change. This supports one of our earlier explanations of the activation energy curve in Fig. 1, where we have stated that

A Dasgupta et aI /Thm Solid Fthns 295 (1997) 37-42 41

Fig 4 Transmission electron micgographs (TEM) and the corresponding diffraction rings (TED) of S1C:H thin films. (a) and (b) TEM and TED of sample prepared at PT = 0.3 Torr, while (c) and (d) TEM and TED of sample prepared at Pv = 0.5 Torr (e) is the TED corresponding to the sample prepared at 0 6 Torr (1 Torr= 133 Pa).

Table 1 TabIe showing the incorporation of boron (B), carbon (C) and hydrogen (H) atoms in SIC:H thin films at different chamber pressures

P'r (Torr) B/St C/Si H/Si

0.3 8×10 -4 2× 10 -9 - 0 5 1 × 10 .3 2.5 × 10 .3 - 0 8 5.4× 10 .3 6 6× 10 .3 8 3× 10 .3

the microcrystal l ine film prepared at 0.3 Torr is inadequately doped compared to one prepared at 0.5 Torr. On the other hand, the amorphous nature of the film prepared at 0.8 Torr hinders the doping efficiency in spite of a higher incorpora- tion of B atoms. This film also contains H / S i ~ 8.3 × 10 .3 while in others, it is below detectable limits. This is a signa- ture of its amorphous nature while others are microcrystall ine with expectedly low H contents (since the crystallites are necessari ly devoid of hydrogen) . Again, as carbon is incor- porated only in the amorphous matrix and grain boundary regions of the film, its content increases with an increase in the fraction of amorphous phase of the films.

5. Conclusions

Highly conductive p-type #c -S iC:H has been prepared by photo-CVD. The chamber pressure plays an important role during deposition of microcrystal l i te thin films. This param- eter also determines the efficiency with which dopant atoms are incorporated. There is an opt imum pressure, depending upon the other deposition parameters, to achieve high con- ducting doped /xc-S iC:H film by controlling crystall inity in conjunction with doping efficiency.

Acknowledgements

Two of the authors (AD and SG) gratefully acknowledge the C.S.I.R, India, for extending financial support.

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