transport properties of doped silicon oxycarbide microcrystalline films produced by spatial...

ELSEVIER Solar Energy Materials and Solar Cells 41/42 (1996) 493-517

Solar F.ner~ Male~ls and Solar Cells

Transport properties of doped silicon oxycarbide microcrystalline films produced by spatial separation

techniques

R o d r i g o Mar t ins a, M a n u e l a Vie i ra b, Isabel F e r r e i r a a, Elvira F o r t u n a t o a, L e o p o l d o Gu imar~es a

a Materials Science Department, Faculty of Science and Technology, New University of Lisbon, P-2825 Monte de Caparica, Portugal

b CEMOP/UNINOVA, Quinta da Torre, P-2825 Monte de Caparica, Portugal

Abstract

This paper presents results of the role of the oxygen partial pressure used during the deposition process on the transport properties exhibited by doped microcrystalline silicon oxycarbide films produced by a Two Consecutive Decomposition and Deposition Chamber system, where a spatial separation between the plasma and the growth regions is achieved. This paper also presents the interpretative models of the optoelectronic behaviour observed in these films (highly conductive and transparent with suitable properties for optoelectronic applications) as well as the interpretation of the growth process that leads to film's microcrystallization.

I. Introduct ion

Wide band gap microcrystalline silicon (~c-Si) carbide base alloys have aroused considerable interest since they combine some advantages of amorphous and crystalline materials [1]. However, these materials can not be produced by conventional r.f. glow discharge process since this technique does not allow the production of /zc-Si films with carbon (C) content above 10 at% [2]. In previous work [3,4], high quality silicon carbide materials were produced, using a strong beam of excited hydrogen [H] to decompose the silicon-and-other mixed-containing gases. This is the case of producing doped microcrystalline silicon oxycarbide films (/.~c-Six:Cy:Oz:H) presenting a high chemical stability, optical gaps (Eop) , and conductivities (o- a) that fit the requirements for solar cell's applications as doping contacts (decrease of the front layer optical losses and enhancement of the optical

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494 R. Martins et al. / Solar Energy Materials and Solar Cells 41/42 (1996) 493-517

path, when used as the back doping layer together with a high reflective metal back contact).

In this work we present results concerning the parameters that control the film's microcrystallization during the deposition process in a Two Consecutive Decompo- sition and Deposition Chamber (TCDDC) system [5], emphasising the role of oxygen partial pressure ( P o ) on the electro-optical properties exhibit by the /xc-Six:Cy:Oz:H films produced. Besides that, the interpretative models of the growth process and of the electrical conduction mechanism, able to explain the behaviours observed, are also presented.

2. Experimental procedure

The films were prepared in a TCDDC system whose main characteristics as described elsewhere [5], are the following: (i) the dilution/reactant gas (Hydrogen, in this case) is directly introduced in the region where plasma ignition takes place; (ii) the gas or mixed gases containing the species to be deposited are introduced close to the region where the growth process takes place (heated substrate); (iii) a bias grid is located between the plasma ignition and the growth regions, leading so to the spatial separation between the chemistry of the plasma and that one of the growth region. This grid also controls the type of species impinging on the growth surface, mainly the charged ones. When a partial base pressure is used, all the system is evacuated to pressures of about 10 -7 mbar after which a back pressure containing the desired gas is established, over all the deposition chamber. In the present case the partial back pressure was chosen to be of Oxygen ( P o ) to observe its role on improving the materials optical gap without decreasing significantly the materials electrical conductivity O'd, as required in doped materials for solar cells applications.

The choice of the deposition parameters for producing the doped /zc- Six:Cy:Oz:H films were anticipated by a study to determine the hydrogen (H) dilution rate and power densities (dp) that leads to film's microcrystallization, as well as to infer the main deposition parameters that control the film's growth rates (GR). To do so, dp was varied from 0 to 150 mW cm -3, for H dilution rates higher than 97%, keeping T s = 250°C and the gas flow, Q = 150 sccm. Once defined the deposition conditions for film's microcrystallization, the H dilution rate, the methane to silane gas mixture ratio (defined as X = CH4/SiH 4 ratio) and dp were fixed (see Table 1), varying only Po2"

The film's structure, composition and morphology were analysed by means of grazing X-ray diffraction, Electron Secondary Chemical Analysis (ESCA), Sec- ondary Ion Mass Spectroscopy (SIMS) and Fourier Transform Infrared spectrometer (FTIR).

The above measurements were performed on films deposited on crystalline silicon and alkali free glass substrates, being the first ones used to perform films' compositional analysis while the second ones were used to perform film's structure and electro-optical transport analysis.

R. Martins et al. / Solar Energy Materials and Solar Cells 4 1 / 4 2 (1996) 493-517

Table 1 Summary of the deposition conditions used, where Po2 varied from 10 -3 to 10 "1 mbar

495

~c n-type /xc p-type

Substrate temp., T s (°C) 220 220 Deposition pressure, Pd (mbar) 0.5 0.5 Gas flow (sccm) 100 100 Gas doping level (at.%) 0.5 2.0 Power, dp (roW/era 3) 100-100 120-135 H 2 dilution ratio (at.%) 95-90 90-85 CH4/SiH4, X (at.%) 25-40 25-40

The films' optoelectronic analysis were performed under vacuum conditions (by a cryostat or vacuum chamber) being the produced samples kept in a vacuum container before their installation on the corresponding measuring chamber system. The film's electrical properties were investigated through the dependence of the static ~r d on the temperature (T) in the range from 350K to 80K, computer controlled, allowing a time interval between measurements of about 2 min, for the complete film's relaxation. Besides that, the time dependence of ~r d at T = 80K, was investigated to determine the role of the a-tissue (that controls the carriers' transient conduction behaviour, between grains) on the electrical transport behaviour. The film's optical properties were determined by measuring the optical absorption [a(A,T)] (with T ranging from 300 to 80K and the wavelength, ,~, ranging from 300 nm to 1500 nm, respectively) using an U V / N I R spectrometer, equipped with an integrated sphere and a cryostat, computer controlled.

2.1. Role o f the deposition conditions on film's microcrystallization

Prior to produce the/xc-Six:Cy:O z H films, an extensive work was performed to know the deposition conditions that lead to the film's microcrystallization. As it is now well established, the main responsible for film's microcrystallization is the presence of a reactant gas (like H), complemented by the utilisation of high dp and /o r T s, during the deposition process [6-9], to allow the growth mechanism to reach the so called quasi chemical equilibrium [10]. Under these conditions, the main role of the reactant gas on the set of chemical reactions that take place on the growth surface is to favour the reconstruction of Si-Si bonds, if the radicals formed during the plasma process have enough energy to eliminate the presence of weak bonds on the growth surface, through an etching process [11]. To enhance the set of reactions on the growth surface, besides using a high dp we have also tried the presence of an ultra-violet light assisting the deposition process [9]. Nevertheless, the obtained results indicate that the main parameter that influence strongly the microcrystallization process is the percentage of the reactant gas (H) present and the d o used. Considering only the role of do, w e observe that the transition from the amorphous to the microcrystallization deposition conditions, corresponds to a high d o to which is ascribed a large drop of the growth rate G R


!0 2

10 ° >-.,

10 -2

¢ 10 -4

10 -6 0

0.6

0.4

0.2

50 100 dp (mW/cm 3)

150

5

4

3-2

2

1

0 50 100 150

dp (mW/crn 3 )

Fig. 1. Dependence of tr a (a) and AE (b) on d o, for n- (zx) and p-type (O) films, using a H/Sill 4 ratio larger than 97%. Fig. 2b also shows G R for the p-type (O) films.

(more than one order of magnitude), accomplished by an increase on o- d and a decrease on the corresponding activation energy (AE), as can be seen in Figs. 1 (a) and (b).

To explain the experimental G R values recorded, we proposed a model [9] where G R is dependent on species formed in the plasma region that impinge on the growth surface (Fn, the gas dissociation factor (ag), the number of Si atoms per unit volume formed in the film (Nsi), the plasma discharge efficiency (~7) (determines the number of secondary electrons produced by each positive ion and the number of ion pairs per electron per Volt potential difference [12]), the species collision cross section (o-s), the species residence time ( r r) and the density of ion species ( f ) assisting the growth process:

a~F~ 1

GR = Ns----7 1 + (rlmrJ)-l' (1)

R. Martins et al. / Solar Energy Materials and Solar Cells 41/42 (1996) 493-517 497

where the quasi chemical equilibrium condition is reached when:

( % > 1.5 × 10 9 exp - Rg-"-'~ and o-s% > 1.7 x exp - ~ ], (2)

where A G is the Gibb's free energy and Rg is the ideal gas constant. Under this condition, it is expected that the growth rate of ~c-Six:Cy:O~:H

follows the relation:

Fo G R ~- agrlmaXNs i . (3)

AS F n --- 1/4 n M Q / A [9], where n M is the number of gas species available within the system (function of the deposition pressure used, Po), Q is the gas flow and A is the growth surface cross section where species will impinge. For the deposition conditions used (n M = 1.2 × 1016 at cm-3), we expect Fn = 3 × 1013 at cm-2s -1. Experimentally, we observe that G R varies almost linear as Q or dp increase,

o

within the microcrystallization region, being its lowest value of about 0.1-0.2 A/s , for dp> 80 mW cm -2. thus, following F_,q. (3), we obtain a value for ag r/m~ x = 0.2- 0.40. Estimating a discharge efficiency of about 80%, the above result leads to a gas dissociation efficiency of about 30% to 60%, function of the used dp. To improve this value it is required that T~ has to be enhanced. Supported by this result and the experimental data of the behaviour of tr d and A E on dp, for both types of doped films investigated, we fixed the deposition conditions shown in Table 1, with Po2 varying from 10 -3 to 0.1 mbar, aiming to produce p-and n-doped tzc-Six:Cy:Oz:H films highly transparent and conductive for solar ceils applications.

2.2. Role o f powder formation on f i lms structure

During the deposition process, the self bias voltage Vb(t) induced by the r.f. power was time monitored to detect the powder formation by analysing the disturbances that occur in Vb(t) [13]. Some of this powder moves towards the growth surface, either by a randomly diffusion process or by electrostatic forces [14]. Once there, they can influence the film's nucleation mechanics, either by originating pinholes or by behaving as nucleation centres for the film's growth process. Indeed, if some of these particulates are crystallites, they can work as "seeds", able to induce crystallisation during the growth process. Experimental data indicate that the rate of powder formation increases for Po2 > 5 × 10- 2 mbar. Therefore, it is expected that these particulates might influence the crystallite's nucleation mechanisms along with the presence of the O. However, it is very hard to obtain some clear information about this effect in a continuous deposition process, since we do not know when these particulates will land on the growth surface. Thus, to understand the role of powder formation on the nucleation mechanism, preliminary tests were performed by interrupting the discharge mo- mentarily (for a period below 5 seconds) at the beginning of the deposition


process, for p o 2 > 5 × 10 -2 mbar, to allow some of the particulates to land on the film's surface. The objective was to observe if some of these particulates lead to the production of/~c-Si films where the crystallites are congregate around possible "seeds". The preliminary structural data recorded in films deposited on alkali free glasses (to avoid any possible influence of the substrate on the growth process as it could happen if c-Si substrates were used) by TEM show ~c films with average grain sizes ri, of the same order of magnitude as the ones recorded in films produced on c-Si substrates [14].

3. Analysis of the results

3.1. Structure and morphology

The X-ray diffraction analysis were performed on films deposited on alkali free substrates, showing the presence of Si microcrystals with r i in the range of 100-50,~ as P02 varies from 10 -1 to 10 -3 mbar. The crystal's volume fraction (fc) was inferred through the ratio between the experimental peak intensity recorded on the /.rc-films and that one obtained in a c-Si sample, taken as reference [1]. Figs. 2 (a) and (b) show the dependence of r i and fc on P02, for the p-and n-type films produced. The results indicate that fc and r i decrease as P02 increases.

3.2. Dependence o f f i lm composition on P02

The film's chemical composition was analysed by ESCA (Fig. 3 and SIMS (Figs. 4 (a) and (b)). Changing Po2 from 10 -3 to 10 -1 mbar the film's composition (without taking into account the amount of H incorporated) changes from Si89C704 to Si56C19025 , for n-type films produced with an H 2 dilution ratio of 95% and X = 25%, while it changes from Si80C160 4 to Si45C35020, for n-type films produced with an H 2 dilution ratio of 90% and X = 40%. The p-type films show the same behaviour, but for different dp and H dilution as indicated in Table 1.

The SIMS analysis were performed to determine the total amount of phospho- rous (P), boron (B) and H incorporated into the films [1]. Figs. 4 (a) and (b), represent the SIMS depth profiles corresponding to P, B and H distributions, for two typical microcrystalline p-and n-type films produced at Po2 = 10-3 mbar (Si89C704 and Si80C1604, respectively). These data were compared with the ones obtained in P / B c-Si samples and with the Hydrogen content, C H of a calibrated microcrystalline sample. The P and B concentrations inferred were respectively 1 x 10 21 at.cm -3 ( = 2 % ) and 3.4× 10 20 at.cm -3 (=0.68%). These data also indicate that among 80% to 40% of the gas phase species containing P and B are incorporated in the films, about twice as much as observed in a-Si:H films [16]. The experimental data also reveal that by changing Po2 from 10 -3 to 10 -1 mbar, the rate of doping incorporation in both films is kept almost constant, while C H

R. Martins et al. / Solar Energy Materials and Solar Cells 4 1 / 4 2 (1996) 493-517 499

70

60

~ . 5O o<

40

30

20

70 ,

' " 1 . . . . . . . ' [ . . . . . . '

(a) n- type :::::::: .............

10 .3 10 -2 10 -j

PO2 (mbar)

. . . . . . . . ] , , , , . . . .

(b) p - type 60

60

50

4og

' 3 0

20

' " 1

60 O .

~ ) i " ~ . . 5O ,_,50

o < 40 ~ °

" ' 4 0

30

30 ~ 20

20 ,,, i . . . . . . . . i . . . . . . . 10 .3 10 -2 10 -I

Po2 (mbar)

Fig. 2. r i and fc dependence on Po2 [(a) n-type and (b) p-type samples] for X = 2 5 % (0 , O) and X = 40% (m, D).

lOO-

o

o t=

Si Si Si Si Si Si $i Si Si Si

0 0 0 0 0 0_.~0 0 0 0 _ . - c - c ~ c - - c - c . - - - - e - - - -C- - - -~ - - - - -

O 450 Sputter time (s]

Fig. 3. Example of an ESCA depth Si 2p, O ls and C ls spectra for the/zc-Six:Cy:Oz:H sample.


106

104

..~

102

10 ° 0

106

20 40 60 Sputter time (min)

80

~- 10 4

102

10 ° 0 10 20 30

Sputter time (min)

Fig. 4. SIMS depth profiles recorded in an n-type (a) and p-type (b) samples produced atPo 2 = 10 -3 mbar, with a composition of SisoC1604 and Si89C704, without taking into account H.

changes slightly, being in average of about 5 at% and 10 at%, respectively for the p-and n-type samples analysed.

3.3. Role ofpo2 on the material's optoelectronic properties

3.3.1. Optical characteristic of i~c-Six:Cy:Oz:H films The optical transmittance and reflectance measurements allow the determina-

tion of a(A,T) at room temperature (RT) and for 80K < T < 250K, as well as the film's refractive index (nrf). These measurements also allow the determination of the optical gap (Eop) and its behaviour with T.


106

10 5 '7

. . . . I . . . . F . . . . I . . . . I . . . .

a-Si:H

c-Si (a) n-type 10 3 . . . . I . . . . I . . . . I . . . . I . . . .

1.5 2 2.5 3 3.5 4

E (eV)

10 6 . . . . n . . . . I . . . . v . . . . I = ~

a-Si:H

,, : c-Si '¢ ,"

," O 0 • I / / :

• , • I ,' • • / • 10 4 , ' • • t •

, - " ] c-Si (b) p-type

10 3 1.5 2 2.5 3 3.5 4

g (eV)

_. 10 5

&

Fig. 5. Dependence of a on E for (a) n-type (where the symbols I , O, • and • corresponds to samples prepared at Po2 = 10-3 mba; 10 -2 mba; 5×10 -2 mba and 1.5×10 -1 mba, respectively) and (b) p-type films (where the symbols I , • and • corresponds to samples prepared at Po2 = 10-2 mba; 10-1 mba and 2 × 10- i mba, respectively). The corresponding Po: and Eop of the samples analyzed are shown in Table 2.

Figs. 5 (a) and (b) show the dependence of a(Eph) on the photon energy (E), for p -and n-type /~c-Six:Cy:Oz:H films. There, it is also shown the typical be- hay•our of a-Si:H, c-S• and/.~c-Si films' spectra, to compare them with the ones of the doped p~c-Six:Cy:Oz:H films. The data show that the p-and n-doped /zc- S i x : C / O z : H doped films are more t ransparent than the a-Si:H, /~c-Si and c-S• films, if C + 0 2 content is higher than 30%.


Table 2 Optical gaps of n and p type ixc-Six:Cy:Oz:H films produced with X = 40% at different Po2 and using different fittings

Sample p o 2 E04 Eop (Tauc 's plot) E2/3 (mbar) (eV) (eV) (eV)

n -type 10-3 2.14 2.15 2.28 10 -2 2.25 2.45 2.60 5 × 10 -2 2.77 2.57 2.76 1.5 × 10 -1 2.72 2.71 2.89

p-type 10-2 2.10 2.17 2.47 10-1 2.35 2.62 2.70 2 × 10 -1 2.95 3.15 3.24

The energy for which a(E)~-10 4 cm -1 (E04) supply information about the degree of film's transparency and it is common used as a good approximation for Eop [16]. The E04 values for the samples shown in Figs. 5 (a) and (b) are enhanced either by changing X or Po2, during the deposition process (see Table 2). The experimental data also indicate that for a ( E ) < 10 4 cm -1, the recorded spectra have a shape similar to the that one exhibit by the c-Si films while in the high absorption regime tends to be more likely that one exhibit by the a-tissue.

The most common method to determine the Eop of a-Si:H base films is by the Tauc's plot [17]. Nevertheless, the/zc-Six:Cy:Oz:H films is a composite, where the composition of the a-tissue is different from that one of the crystallites and so, it is expected that a(E) shows a behaviour as [18]:

a ( E ) = (E - Eop) n, (4)

where Fagen [18] proposed n = 2/3 in opposition to Tauc's plot where n = 1/2. In these conditions one considers that both types of non-direct and non-indirect interband transitions (without the involvement of phonons assisting the process, to conserve energy) can occur [18].

Fig. 6a-6d compare the Tauc's and the [ot(E)E] 2/3 plots, for the n-and p-type films of Figs. 5 (a) and (b). The data show that in the high absorption region (a(E) >_ 104 cm -1) both types of plots exhibit straight lines behaviours. Besides that, it is observed that the recorded optical gaps are influenced by Po2, as shown in Table 2. Under these conditions, we expect that the optical behaviour of /zc-Six:Cy:Oz:H films are mainly governed by the a-tissue. Here, it is important to notice that [a(E)E] 2/3 plots evidences the possible presence of a non-direct to a non-indirect gap, since the low absorption and the high absorption regimes are fitted separately by two linear plots. This result proofs to be difficult to determine Eop in multicomponent films by the conventional theories and so, the Eop obtained can be only considered as a qualitative information about the degree of film's transparency. Fig. 7 shows the dependence of a(T) on E, using as parameter T, for n-and p-type /zc-Six:Cy:Oz:H films produced with X = 25% and Po2 --" 10-3 mbar, respectively. The p-type films reveal an enhancement of or(T) as T de-


t ~

1400

1000

. . . . I . . . . I . . . . I . . . . I . . . .

(a ) . 2 .15 eV

• 2.45 eV • 2.71 eV • 2.57 eV J

600 b

i " "

I I • •

200 • n-type

1 . 5 2.0 2.5 3.0 3.5 4.0

E (eV)

. . . . i . . . . t . . . . I . . . . I . . . .

{b} ,, 2.28 eV

• 2,60 eV

• 2.89 eV •

• 2.76 eV

; •

• O •

2000 • • n-type~

1 . 5 2.0 2.5 3.0 3.5 4.0 E (eV)

14000

10000

6000

1400 . . . . I . . . . t . . . . I . . . . .

(C) • 2.62 eV

3.15 eV 2.17 eV

1000

f / . . 600 ~ ~ .

L 200 p-type

10000

6000

2000

2.0 2.5 3.0 3.5 4.0

E (eV)

{d) • 2.70 eV

1 • 2.47 eV

• 3.24 eV :

" I

: ' , : .

/ g " •

p-type i i i , I , i i i I , i , i ] _ , : L ,

2.0 2.5 3,0 3.5 4.0

g (eV)

b) d)

Fig. 6. Dependence of (aE) 1/2 and (orE) 2/3 on E, for: (a),(c) n-type and (b), (d) p-type films, using different Po2 as indicated in Table 5.

creases, while the n-type samples show an opposite behaviour. The role of T on Eop was also determined by following a similar procedure as that one proposed by Cody [22]. The samples analysed have thicknesses below 0.5 Ixm.

The insert in Fig. 7 shows the experimental recorded Eop obtained as proposed by Cody [19], where Eop(T) fits an Einstein oscillator [20], given by:

A E°p(T) = e°do) - (5)

e T - 1

5 0 4 R. Martins et at/Solar Energy Materials and Solar Cells 41/42 (1996) 493-517

'9

1 0 7

1 0 6

' ' ' I

2 . 1 5

. - , 2 ' 1 0 I

f . o 2 . 0 0

1.95

1 .90 t

' ' I ' ' ' I . . . . . I ' '

o n - t y p e []

[] p - t y p e

- - s i m u l a t i o n

, ~ - , , ! ! t ! ings, . . . . . . . . . . . . . .

5 0 100 150 2 0 0 2 5 0 3 0 0 3 5 0 ~ . -

T ( K )

p-type , 1 105 ~____ . . . . ~ , "

....~ ~ T=80K

. . - c ~ - T = I 5 1 N

L 2 ~ T = 2 0 5 K n-type , ,, = ' ~

. . . . S : , / ~ j . . . . T=250K . . - - - T = 2 5 6 K 1 0 4 , , , I , , , I , , , P , , , I , , , I , , ,

2.0 2.2 2.4 2.6 2.8 3.0 3.2

E (eV)

F i g . 7. a ( T ) spectra for n- and p-type samples prepared with X = 2 5 % a n d X 4 0 % , r e s p e c t i v e l y . P o 2

was 0.05 mbar and the values of T are labelled in the figure. The spectra corresponding to the p - t y p e

sample was shift upwards by a factor of 10, for clarification reasons. The insert shows the Eop(T), for both types of films.

where for p-type samples, the signal in the second term of the above equation is positive, showing an abnormal behaviour from what is expected, attributed to an enhancement on films' internal static disorder and so, to an enhancement of the bulk density of states, caused by the presence of boron atoms, as also suggested by the A and 0 e values recorded, respectively of about 0.2 eV and 400K and 0.42 eV and 445K, for n-and p-type samples. As the measurements of Eop(T) have been taken below room temperature, the role of thermal disorder on the values recorded can be neglected, as observed by Weiser et al. [21] in a-Si:H and c-Si films. This result agrees with the previous data [15], where it is observe that the p-type films are more defective than the n-type ones. The data recorded also lead to Eop(O) values of about 2 eV, showing that most of the optical transitions are governed by the a-tissue [22].

The volume fraction of voids (fv) can be estimated by knowing the value of the refractive index n o near the spectrum infra-red region, once it reflects the degree of the tissue compactness. So, its variation from the expected value for the ordered


45 n-type . ~ ~ ..... . _~-- ~ . . . . . .

~ 35 ,

25 ,/

~' (a) --o--- f; (X=25%) 15 r . . . . . . I . . . . .

10 .3 10 .2 10 -J

P02 (mbar)

45 p-type . i i

35

,,.~>

25

(b) + f~ (X=25%) 15 . . . . . . . . I . . . . . . . 10 .3 10 -2 10 -I

Po2 (mbar) Fig. 8. fv dependence on Po2 [(a) n-type and (b) p-type samples] for X = 25% and X = 40%, as labelled.

c-Si films, it leads to infer fv. This type of study was pe r fo rmed by Gode t [23] that obta ined the following expression for fv:

( N Z - n 2 ) ( 1 + 2n 2)

f~ = 3 n Z ( U 2 - 1) ' (6)

with N(Si ) = 3.458.

Figs. 8 (a) and (b) show the dependence of fv on Po~, for both films analysed, ei ther for X = 25% or X = 40%. The data show that both type of films present similar behaviours, being fv smaller when X = 25% than when X = 40%.

3.3.2. Behaviour o f ad and the role o f the oxygen partial pressure The tr d measurements were pe r fo rmed in the t empera tu re range of 300 to 80K

allowing the de terminat ion of tr d at R T and the cor responding activation energy, A E. In Fig. 9a we depic ted the Arrhenius plot for both types of films analysed. The behaviour observed shows that in both types of films the experimental points do

506 R. Martins et a t / S o l a r Energy Materials and Solar Cells 4 1 / 4 2 (1996) 493-517

101

10 °

]0-~

10 .2

10 -3

' I ' ' ' 1 ' ' ' 1 ' ' ' 1 ' ' ' l ' E ' l ' ' ' n - t y p e

- - p-type ~P I

1.96 eV-

- - " . . . . . . . - 2.30 eV

2.33 eV

2.40 e V

~ 2 5 4 e V ! (a)

2 4 6 8 10 12 14 16

1000/T (K -I) 101

I ~ n ' ( y p e " ~ ' " ' " ' ' " E ' ' '

100 ~ 1.96 eV

~ 2.3eV

10 -I 2.33 eV

2.40 eV

10 -2

(b) 2.54 eV

10-3 , , ~ , , , i . , , i , , , i , , , i , , , i , , 0.24 0.28 0.32 0.36

T-t/4 (K -t/4 )

Fig. 9. (a) o" d versus T - 1 for n-and p-type films produced using X = 40% at different p o 2 (see Table 2) that lead to Eoo labelled in the figure. (b) Dependence of ~r d on T -1/4 for the same films as in (a).

not follow the expected tr a versus T -1 linear plot. In Fig. 9b we represent the same tr a but now for a T-1/4 dependence, where a good fitting of the experimental points is obtained, specially for films exhibiting a high transparency (Eop > 2.2 e V) and conductivity (o'd > 10-1 O-lcm-X), produced with X > 40%. This result indicates that the electrical conduction mechanism is mainly performed by percolation between grains where the intergranular tissue plays an important role in how the percolation channels are formed [24], as expected from the percolation model as will be described below. The above data also indicate that the number of phonons available to assist the conduction process at low T is very low and so, the carriers are " t rapped" in localised states from which they can travel, by hopping, to the closest available site, function of the state's capability in retaining (or not)

R. Martins et al. // Solar Energy Materials and Solar Cells 41/42 (1996) 493-517 507

Fig. 10. Two-dimensional "Swiss cheese" model of random continua.

the carriers, which explains the tr d time (t) dependence behaviour observed at T = 80K, where for the n-and p-type samples studied o- d follows a behaviour, given by:

~a(T , t ) = t r a ( T ) e x p ( - tr ), (7)

where ~d(T) is the conductivity at T, for t = 0 and t r is the relaxation time whose value is dependent on film's composition (for an n-type film with Eop = 2.5 eV, t r = 476 min. and for a p-type film with Eop --- 2.3 eli, t r = 625 min.). The high t r shows that the conduction, though between grains, it is also dependent on the amorphous encapsulant, where the percolation channels are established, like in a "Swiss cheese" [25] (see Fig. 10). That is, the a-tissue time modulate the tr d behaviour, mainly function of the existing shallow states in the grain boundaries and on the bulk of the a-tissue.

4. Conduction mechanism in doped itc-Six:Cy:Oz:H films: proposed electrical conduction model

The existing models to explain the conductivity, o- d behaviour of /xc-Si base films, range from the classical ones (based on the counterpart a-Si tissue) to those based on the conduction behaviour observed in polycrystalline silicon. Thus, as the /xc-films are essentially constituted by small grains spread in the a-tissue, it is expected different transport behaviours for both components. Indeed, the carriers flow is dependent on grain's conductivity, its spatial separation and on the nature of the intergranular material. As a first approach, /xc-fllms can be considered as "composites" of homogeneous regions in which the conduction mechanism is ohmic, as in a single crystal, if the crystal clusters are sufficiently large. However, if r i is small enough, it is possible that the carriers' mean free path to be of the same order of magnitude than r i. This behaviour is explained by two main models. One

508 R. Martins et al. / Solar Energy Materials and Solar Cells 4 1 / 4 2 (1996) 493-517

is based in the hopping-like conduction mechanism, observed by the first time in 1950 by Hung et al. in Ge [26]. This conduction mechanism was further studied by Mott [27] and Conwell [28]. This initial model was further developed by Miller and Abrahams [29], considering the approach of the structure to a discrete random resistor network model (the so-called reduced resistive model), to interpret the carrier's flow on structures of a granular type.

As far polycrystalline materials are concerned, the standard treatment of the electrical transport properties is based in the paper presented by Petritz [30], where he showed the dependence of the film's resistivity on the barriers formed between grains. More recently, Snejdar et al. [31] and Jerhot et al. [32] have considered a generalised model that includes the potential barrier between the grains and the inter-granular material.

One of the main criticism about the existing models is that they consider uniform sized grains and uniformly thin inter-granular material without considering the probability of forming conductive chains and/or the probability of the excited carriers to hop or tunnelling between grains. These deficiencies were corrected by the work of Ambegaokar et al. [33], Shklovskii et al. [34], Pollak [35], and Adler et al. [36], that explain the conduction mechanism on a "disordered" network by the percolation theory, assuming a continuous network where conduction sites can be either occupied or empty, leading to a structure model that resembles a "Swiss cheese" with channels between "grains" through which carriers can percolate, as shown in Fig. 10.

In the classical percolation problem, regions of different conductivities are assumed to randomly fill the space and the conductivity is calculated as a function of the volume fraction of each component. If there are two components where one of them is conducting and the other not, then there is a discontinuous transition between an overall conducting and an overall non conducting state, when more than 10% of the total volume is occupied by the conducting material. The transition point is known as percolation threshold (Pc). Under this condition, ¢r d is dominated by the existing barriers between grains even though, the total volume fraction occupied by the barrier region is low.

Under these conditions, the overall ~r d is calculated as a function of the fraction of conducting bonds and of the film's composition [37] which tends to be zero below a critical percolation threshold value, Pc. Generally, it is observed that ¢r d follows a power law behaviour given by [37]:

= 0(p - p c ) ' , (8) where t = 1.65 in 3-dimensions and t = 1.1 in 2-dimensions case [38].

In the following we discuss the percolation model assumed to be the most probable conduction mechanism in multicomponent/zc-films above analysed.

4.1. The percolation model

To explain the conduction mechanism by percolation, let we consider the system as constituted by several interconnecting resistors between sites and a distributed


,a,[

(b) %

p

i

J

b) Fig. 11. (a) Schematic picture for the links (one dimensional chains), nodes (crossing points of the links) and blobs (dense regions with more than one connection between two points, shown as circular, that correspond to grains) of the infinite cluster slightly above the threshold. The thin lines are the dead ends (for clarity only few are shown [24]). (b) proposed infinitesimal discrete electrical model shown the correlation between two adjacent grains in the microcrystalline structure.

capacitance that controls the carriers transfer between inter-sites (adjacent or randomly distributed but close enough to allow the carriers to percolate), dependent on the characteristics of the a-tissue (see Fig. 11). We also assume that the elemental components will tend in the limit to form a continuous network, behaving as a "Swiss cheese-like" structure with established channels between grains through which carriers can percolate. Considering that: 1. the existence of a net of isolated clusters of high conductivity, such Rij < Rc,

where Rii is the resistance between sites i and j and R c is the critical percolation resistance; small number of resistors with R o. = Rc which connected together make a subset of high conductance clusters, forming a sample-spanning cluster;

2. remaining resistors with Rq >> R~; 3. the equivalent occupancy probability of site i is represented by a capacitance C i

e 2

(C i = ~ a T e x p ( - E i / k B T ) , where e is the electron charge and E i is the energy

of site i, measured from EF), considering a cross it the existence of a potential z/.

4. the carrier's transition probability between sites (i and j, for instance) is governed by the a-tissue, represented by an equivalent capacitance Cq;

5. all potentials in the model has to be referred to the a ground potential;

510 R. Martins et aL / Solar Energy Materials and Solar Cells 41/42 (1996) 493-517

and if an electric field ff is applied, the current flow, in unsteady conditions, follows the relationship (in agreement with the Kirchhofffs rules [39] and the energy conservation principle) [29,36]:

CiAV i avij Z A-~ov~Cij-~ "~j(Vj- Vi)o'ij, (9)

where the network is considered as composed by a set of links-nodes-"blobs", as proposed by Stanley [40].

In the ohmic regime, the current is proportional to ff and it is dependent on the probability of a site being or not occupied (where trij is the normalised conductivity, being ~ j = 0 when p <Pc and gii = 1 otherwise) and on (V: - V,.), the differen- tial potential between 2 sites.

In the quasi-steady state condition this probability can be given by the Fermi distribution:

1

exp / ~--~a'T) + 1

where the superscript refers to the quasi-steady state condition. Considering that V i ct f t . ~i and V: - V/ct f t . ?ij (rij is the distance between sites

i and j), the time dependent term on the left side in Eq. (9) can be related to the site frequency occupancy and on the network relaxation behaviour.

In quasi-steady state we have C(o,ij)) = [pO eZ(1 _ PiO)]/(kB T) and so, Eq. (9) is re-written as:

"~i ei°e2(1 -- eiO) OVi kBT ~t = j t \ V i - - f f ' r i z - - ( I1 (11)

where Pi is such that Pi = Pi ° + APi;

[ ( E i - e V i ) ] -1 Pi=Pi° + APi = exp kBT + 1 (12)

From the above equation we define an impedence Zq between any two junctions i and j such that:

(zij)-l= ~Br(,0)-lexp ~ +2 , (13)

where a is the spatial distance through which carriers localisation occurs, dependent on the Bohr radius [36] and r o is the average time for carriers to flow from site i to site j.

The total impedance (Z) is obtained by integrating the impedance between sites i and j (Zq) over the entire bulk of the material. However, as the carrier's conduction is dependent on the easiest formed path where carriers require the lowest consumption of energy, we can estimate the required average energy

R. Martins et al . / Solar Energy Materials and Solar Cells 41/42 (1996) 493-517 511

considering that exists a critical estimate conductance (G c) to which is ascribed a Zc such that:

2 Zc wC c , (14)

that corresponds to the largest value of the conductance such that an element of the network with Gij > G c still contains a conducting sample-spanning cluster and to a frequency at which a sample-spanning cluster is formed by the first time.

From the above statements, the most probable condition to happen is Z c = Rc, where one can not neglect the exponential dependence of R c on the site energies. This means that conduction is not performed through the nearest neighbour site but through the one with the consumption of the lowest energy, by the carriers involved. Thus, to determine Z c we must know the value of this critical energy that determines the number of carriers involved in the process, per unit volume (n). To do so, let we consider that the density of states g(EF), where carriers are spatially localised (clusters), can be considered constant within an interval of a few kBT:

n = g ( E F ) ~ c k B T (15)

with ~c = E i j / ( k B T ) + 2ri j /a being a random variable, dependent on the energy correlation between two sites in which percolation can occur and on the distance between the corresponding sites.

Under these conditions, the probability to find available a site is given by:

1 g( E F ) k n T = ( V ) ' (16)

where ( V ) is the average volume, assumed to be:

4 ( V ) = ~ r r 3 . . (17)

Now, if each site can be approached to a sphere, we get:

rij = -~ ku T , (18)

and so:

4 3 ~" 3( ~:c Ei )3. (19) ( V ) = ~Trri j = ~ a 2 kBT

The critical Z~ is then determined by knowing the maximum of the exponent in Eq. (13) for which r ~ 0. This is accomplished by substituting the average value of rij as a function of E i in Eq. (13) and taking into account the dependence of ~:c on Eij. This leads to the determination of the minimum energy required for carriers to percolate between grains. By doing so, we obtain the following equation (for a 3-dimension case):

o o0ex [

512 R. Martins et aL / Solar Energy Materials and Solar Cells 41/42 (1996) 493-517

where T oct 1 6 a 3 / ( k B T ) with a' being the spatial attenuation factor, dependent on g < E F) and G o is the pre-exponential factor of the conductance, dependent on sites, number of channels available and film's composition [33].

Eq. (20) indicates that when the electrical conduction % in a given structure is due to percolation, it is expected a temperature behaviour such that tr d ~ T -1/4, where phonons can assist the conduction process, not implying necessarily that the carriers transport is due to variable range hopping around EF, in agreement with what was also proposed by Adler et al. [36]. This type of model fits well the electrical conduction behaviour observed in the samples analysed in this work and so, it can be considered that the most probable conduction mechanism in microcrystalline films based in multicomponent or composite like structures is by percolation.

5. Discussion and summary of the results

The obtained results show that the presence of O in the deposition process controls CH, ri, fc, and fv, enhancing also the chemical reactivity on the growth surface and so, C incorporation. However, as P02 varies, the rate of O and H incorporation in the n-and p-type films are different. The O is more incorporated in the bulk for the n-type than for the p-type films, approaching a 1:1 ratio between the 0 2 present in the gas phase and O incorporated in the film as P02 increases, whilst the H bonded to Si is higher (for the same X ratios) for p-than n-type films.

The ESCA data also reveal that the amount of O incorporated in the films as suboxides decreases asp02 increases, mainly when P02 > 10-2 mbar.

ri, fc and fv are also dependent on Po2 (up to Po2 = 5 × 10 -2 mbar, above which Po. enhances the degree of film's disorder) and on the substrate surface structure ~1]. Figs. 12 (a) and (b) show the role of Po2 on the optoelectronic properties of the doped tzc-Six:Cy:Oz:H films analysed. There, it is plot the dependence of the Eop and cr a at room temperature on Po2, for the set of n-and p-type samples investigated. When comparing these results with the ones obtained for silicon carbide n-and p-type/xc-films produced by ECR [4] (where the n-and p-type doped films have two orders of magnitude difference in o-a), we observe that films produced by the TCDDC technique show similar behaviours for both n-and p-type films.

The experimental data also show that Eop depends on X used and it is directly proportional to Po2 while % shows a reverse dependence on Po2" For Po~ > 5 × 10-2 mbar, the optoelectronic properties tend to be more dependent on O than on C content. Thus, for Po~ < 5 × 10- 2 mbar, O compensates mainly vacancies, acting as donor like states ('similarly to what happens in degenerated oxide based semiconductors [20], not favouring the appearance of O - O bonds.

When Po~ > 5 × 10 -2 mbar the O - O bonds are favoured, leading to an increase of SiO 2 clusters, as observed by ESCA and confirmed by IR data [15]. This enhancement on C and O content leads to an enhancement on Eop. On the other


>

O

, 8 i i ) i i i t l I r i , , , , , , [

2 . 6 ~ •

2.4

2.2

2 (a) n-type

1.8 . . . . . . . . ' . . . . . . . . a 10 -3 10-2 10 =1

Po2(mbar )

2.8 . . . . . . . . ~ . . . . . . . . )

10 t

1 0 ° ~ "

101t~ "~

10 2

l 0 t

2.6

> 2.4

o 2.2

(b) p-type 1.8 . . . . . . . . I . . . . . . . . I

10 -3 10-2 10-1'

Poz (mbar )

100

l 0 t tz "~

16 2

Fig. 12. Eop and ~r d dependence on Po2 for p- and n-type [(a) and (b)] films produced at X = 25% (0 for Eop and O for tr d) and X = 40% (e for Eop and O for trd). The H dilution ratio and the dp used are listed in Table 1.

hand, the fast decrease in tr d for Po2 > 5 >( 10 -2 mbar can be related to the enhancemen t o f fv (associated to the decrease of r i and fc) that leads to the format ion of a more porous material and so, to a more disordered structure. The t empera tu re shift o f Eop(T) on T follows the expected Tauc ' s Eop red shift [19]. However , for n- type samples, the shift leads to the shrinkage of Eop as T increases (C = 0.2 eV and 0 E ~ 400K),as observed in a-Si :H films. Forp- type films the shift widens the Eop as T increases (C ~ 0.42 eV and 0 e ~ 445K), similarly as observed in materials with a high D O S (for instance, c-Si shows a 0 e = 350 K, that cor responds to an oscillator energy of about 30 m e V [20] that increases as the film's disorder increases [22]. This indicates that the degree of disorder is higher for p- type than for n- type films, confirming the experimental da ta recorded by IR

514 R. Martins et al. / Solar Energy Materials and Solar Cells 41 / 42 (1996) 493-517

Table 3 Characteristics of doped/.~c-Si alloys, from different authors

Alloy Eop ~rd AE r i fc Reference (eV) (flcm) -1 (eV) (.~) (%)

Si:CI:H p-type 2.20 2.5 0.08 75 [48] n-type 1.95 4.1 × 10 .2 0.22 50

Si:N:H p-type 1.95 5 × 10 -3 0.1 100 30 [49] n-type

Si:F:H p-type 1.30 0.6 0.02 73 71 [50] n-type 1.50 4.5 0.02 98 30

Si:C:H p-type 2.08 10 -3 [51] n-type 2.00 5 )< 10 -2

tzc-SiO p-type 2.18 1 × 10 -2 [52]

and PDS [1,15] as well as the behaviour observed for the tra(T), function of the used Po2"

Considering the set of results published for/zc-Si base materials, in Tables 3 and 4 we present the main characteristics related to Eop and trd, type of structure and the deposition technique used, for doped films with different origins. From these tables it is concluded that the/zc-Six:Cy:Oz:H films are a potential candi- date to be used as a dopant layer in a variety of devices where wide band gap materials, highly conductive and transparent are required, such as solar cells (enhancement of the blue response of the device and the light optical path when used as doping layers [41] or even in the production of a complete microcrystalline solar cells as already tried with /zc-Si [42,43] or in the production of thin film transistors [44-47], taking advantage of the optoelectronic characteristics presented by this material.

The data also show that the preparation technique, a spatial decoupling of the decomposition and deposition areas, leads to the formation of/xc-Si phase even at high 0 and C incorporation ratios. This is in contrast to conventional r.f. glow

Table 4 Electro-optical characteristics of/zc-Si alloys produced using different techniques

Depos. Alloy Material Eop ~rd A E Reference system (eV) (Ocm) -1 (eV)

ECR Sil_x:Cx:H p-type 2.4 0.1 0.1 [53] n -type 2.3 200

VHF-GD Si:H/Sil.~:Cx:H p-type 2 .06 0.8/0.1 0.01/0.08 [54] n -type

TCDDC-GD Six:Cy:Oz:H p-type 2.44 2 0.020 [1] n-type 2.3 1.2 0.032

Photo-CVD Si:H p-type 2.3 1 [55] n-type 2.0 21 0.02-0.03

GD Si:H/Sil_x:Cx:H p-type 2.08 10 -3 [56] n-type 2.0 5 × 10-2


discharge methods where the crystalline phase does not appear for concentrations typically higher than 10 at% of C in the growing film [2].

The films under analysis do not show the presence of Si:C microcrystals, in contrast to those ones produced by ECR [4], with similar optoelectronic properties. One of the possible reasons for this disagreement is the presence of a grid in the TCDDC system (between the plasma and the growth regions) that slows down the energy of the H species before reaching the growth region.

The results obtained also reveal that the presence of O in small amounts Po2 < 5 x 10 -2 mbar), in the deposition process, enhances the doping incorporation and the film's crystallisation, leading to the production of wide band gap films, highly conductives. Nevertheless, still remains work to be performed, mainly concerning the passivation of the surface voids related to the columnar like structure of the material, to avoid micro-shunt paths and to be able to produce noise free and ohmic-like contacts. It is also important to notice that the transport properties are well interpreted by the two phase models proposed, that associates the electrical conduction to the /zc-islands (controlled by the a-tissue) and the optical transitions in the high energy regime to the a-tissue and to the fact that we are in presence of a multicomponent material.

Acknowledgements

The authors are grateful to A. Ma~arico and M. Santos in the help given during film's preparation as well as to G. Willeke for the inestimable contribution during the first phase of this study. We also would like to thank F. Lopez and J. Puigdollers of the University of Barcelona for performing the SIMS analysis. This work was supported by NATO SfS programme (PO-THINFILM project), by Junta Nacional de Investiga~o Cientifica e Tecnol6gica, project N ° STDR / 6 6 1 / C T M / 9 2 / A and by Programa Estrat6gico de Dinamisa~o e moderniza~o da Indfistria Portuguesa, project N ° 0250. The authors also thanks the co-funding of "Funda§~o Luso Americana para o Desenvolvimento" for the grants given.

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