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Importance of surface processes in defect formation in a-Si:H
Journal of Non-Crystalline Solids 164-166 (1993) 31-36 North-Holland
JOU~NALO~ ~ ' ~ d ~ ~LII~
Section 2. Deposition
I m p o r t a n c e of surface p r o c e s s e s in defect formation in a-Si:H Gautam GANGULY and Akihisa MATSUDA Electrotechnical Laboratory, 1-1-4 Umezono, Tsukuba City, Ibaraki, 305 Japan. The deposition temperature and growth rate dependence of the defect density in a-Si:H is shown to depend on the balance between rates of microscopic dangling bond creation and saturation reactions on the steady state growing surface. The surface dangling bonds are incorporated into the bulk as the film grows. On the basis of this model we show that the defect density can be reduced through precursor assisted defect suppression (PADS) at high substrate temperatures (>350°C), and defect reduction by energized precursors (DREP) at low substrate temperatures (<200°C).
1. INTRODUCTION
2. SURFACE REACTIONS
The properties of PECVD a-Si:H films are a sensitive function of the process conditions, Progress with in-situ characterization of the plasma [1] and the growing surface [2] have allowed us considerable insight into the complex growth process. A link between growth conditions and material properties would allow control of material properties. The defect density (Nd) is the key property controlling the functioning of photovoltaic and other devices made using a-Si:H. Nd is empirically found to vary with substrate temperature (Ts) and the deposition rate (Rd) both of which are of critical importance for device fabrication. In this paper we identify the important microscopic reactions, between the precursors from the plasma and the growing surface that create and saturate dangling bonds (DB). We argue that, during steady state growth, the DB creation and saturation rates result in time independent dangling bond densities on the growing surface, and that these DB are incorporated into the bulk as the film grows. This scheme of defect formation on the growing surface explains quantitatively the experimentally observed Rd and Ts dependences of Nd. Replacing hydrogen with deuterium decreases Nd systematically, confirming a sur-
In-situ diode laser absorption spectroscopy reveals that the number density profile and the consequent flux of -Sill3 in a low pressure silane discharge correlates well with the measured deposition rate given the measured sticking probability of 10% [1]. In-situ infrared absorption spectroscopy (IRAS) of the growing aSi:H surface shows that it is saturated with a monolayer of hydrogen at Ts between room temperature and 480°C [2]. Based on these two results we make the simplifying assumption that growth occurs from -Sill3 arriving on a hydrogen covered growth surface. The reactions of -Sill 3 on the growth surface have been speculated upon previously to explain the measured temperature dependence of the surface reaction probability (13)and sticking probability (s) on the a-Si:H growth surface [3]. The surface hydrogen abstraction reaction of -Sill3 creates a surface DB (Si-) and a silane molecule (Sill4)
face mechanism. Two methods of controlling Nd, at high Ts using PADS, and at low Ts using DREP, are presented.
leads to growth by formation of a Si-Si bond. The recombination reaction between -Sill3
- Sill 3 + Si-H = Si- + Sill 4 ,
(1)
while the DB saturation reaction of -Sill3 -SiH3 + Si- = Si-SiH3,
0022-3093/93/$06.00 © 1993 - Elsevier Science Publishers B.V. All rights reserved.
(2)
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si:H
32
-Sill3 + -Sill3 = Si2H6
(3)
forms a disilane molecule (Si2H6) that can desorb from the surface. In addition to these three reactions of -Sill3, there is the thermal hydrogen desorption reaction Si-H + Si-H = 2Si- + H2,
nated site to another until it reacts. We assume also that when a -Sill3 hops on to a DB site, it reacts. The rate constants of reactions 1 and 2 can be written as [3] Ca = B (1/N o) Avaexp{-Ea/kT} , 1-(N H/No) AVhexp{-E h/kT}
(7a)
C s = B (NH/No 2) AVhexp{-E h/kT},
(7b)
(4)
which creates two surface DB and a hydrogen molecule (H2), and the recombination reaction
1-(NH/No) AVhexp{-E h/kT}
between two DB on the surface that diffuse to suitable locations via migration (tunneling!) of surface bonded hydrogen,
where A and B are normalization constants for the total reaction probability of -Sill 3 on a single site and the total reaction probability of -Sill3
Si- + Si- = Si-Si.
during its lifetime on the surface [3]. Eh is the activation energy for hopping of precursors while Ea is the activation energy for the abstraction reaction 1. Vh and Va are the corresponding attempt frequencies. If we assume Arrhenius type processes for reactions 4 & 5 as
(5)
This reaction creates a Si-Si bond on the surface but does not constitute growth. The reactions 1&4 create while reactions 2&5 saturate DB on the growth surface, but the importance of reaction 3 is because it leads to a flux independent precursor diffusion length, unlike other systems [4]. The -Sill3 concentration on the surface during steady state growth, e = [3~ where ~ is the flux of -Sill3, and x is the reaction lifetime of -Sill3 on the growing surface that is determined by the sum of the rates of reactions 1,2 & 3 [3].
3. GROWTH SURFACE TO BULK DEFECTS We now consider the rates of DB creation and saturation due to the surface reactions outlined above. During steady state growth, the rate of change of the surface DB density dNs/dt = CaNH0 - CsNs0 + CHNH 2 - CdbrNs 2 =0
(6)
where the density of hydrogenated surface sites, NH = No, the total density of surface sites, except for Ns, the orders of magnitude smaller density of DB sites. The C's are the temperature dependent rate constants of reactions 1,2,4 & 5. We assume that when it arrives on the growing surface, a -Sill3 that physisorbs, begins to diffuse by hopping from one hydroge-
CH =aH 2 VH exp{-EH/kT}
and
Cdbr = adbr2 Vdbr exp{-Edbr/kT},
(7c) (7d)
where the a's, v's and E's are characteristic distances, attempt frequencies and activation energies, respectively, then we have in eqn.6, a first order approximation to the steady state DB concentration on the growing surface. The connection between the DB density on the growing surface and the film bulk is established through the hypothesis that the dangling bonds on the steady state growing surface are incorporated into the bulk as the film grows. Then, we have the bulk DB density Nd = (Ns)3/2. The macroscopic growth rate Rd = kCsNs0, where k is a conversion factor. Equation 6 can be extended to include the effects of other precursors, Hradical reactions under microcrystalline silicon formation conditions, impurities and ion bombardment. 3.1. Growth rate dependence Equation 6 is the general equation for obtaining the steady state DB density on the growing surface. However, we do not a priori have values of the activation energies of these different processes. We note that when e is
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si.'H
high enough, the first two terms dominate e q u a tion 6. On the other hand, the third term is important when 0 (Rd) is low and Ts is h i g h . Thermal desorption of surface hydrogen has been observed, by IRAS of the growth surface, to occur only at Ts>350°C. The fourth term of eqn. 6 becomes important when 0 (Rd)is low and Ns (Nd) is high. Therefore, at Ts=250°C, we may neglect the third and fourth terms of eqn. 6 to get Ns = NH Ca/Cs.
(6a)
Equation6aimpliesthatNdshouldbeindependent of 0 and hence, Rd. The variation of Nd, evaluated using CPM, as a function of Rd is shown in fig.1 for a-Si:H samples deposited at
Z
f_ ~ . _ ~kLL{ 15r ...~,~.] . . 1010-i
100
increasingRd. The variation of Nd with Rd at 400°C is shown in fig. 2 for a-Si:H samples deposited from pure silane at a pressure of 30 mTorr by changing the flow rate and rf power.
~
1017 . ~.-. . . . . . . . . ~ ~I • a-Si:HJi il , - ~ : 1160 ~ ~ D . O -
E o
.
~ 1015 '10 ,a-Si:Dq 14 ........ 101(~1 100
~
~
101 '
I
I
•
|===
102
Figure 2. The defect density, Nd vs. the growth rate, Rd for a-Si:H & a-Si:D films at Ts=400°C. 101
102
Rd (/VS) Figure 1. The defect density, Nd vs. the g r o w t h rate, Rd for a-Si:H films at Ts=250°C. 250°C under conditions that strive to m a i n t a i n -Sill3 as the major precursors. This is attempted by keeping the pressure at 30mTorr and adjusting the flow rate of pure silane along with the rf power to avoid formation of higher silane related precursors, which is the likely cause of disparate results in previous studies. Above 350°C, the surface bonded hydrogen desorbs thermally as seen by IRAS, and is the reason for the onset of an increase of R d with increasing Ts, keeping plasma conditions constant[3]. Therefore, the effect of reaction 4 becomes significant compared to reaction 1. Hence, neglecting the first term of eqn. 6 cornpared to the third term, we have Ns = NH2 CH / Cs e.
33
(6b)
Thus, at constant Ts, Nd should decrease with
Distinct from the variation at 250°C we see a clear decrease of Nd with increasing Rd. This can be understood as follows: at a constant temperature,the surface hydrogen is thermally evolved and dangling bonds appear on the surface at a constant rate. These surface DB are saturated by the precursors which arrive on the growth surface. Increasing the rate at which the precursors arrive on the surface suppresses the net density of DB remaining on the surface so that we have precursor assisted defect suppression(PADS)[5]. The data in fig. 2 are in agreement with eqn. 6b provided Rd is proportional to 0. Assumingfurther that NH=No=1015cm2, Ea =0.1eV, Eh=0.4eV, Vh/V a ~10 8, aH2VH/VsXB =10-12cm 2, where "~is independent of Ts & Rd, EH=I.3eV from the measured activation energy of Rd [5], the measured value of ~ = 3.2x1015cm-2s-1 at Rd =l,~s [1] along with the average value of N d at 250°C shown by the dashed line in fig.l, we may obtain fits between Nd and Rd using eqn.6b as shown by the upper
34
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si:H
line in fig. 2. Also shown in fig. 2 are the Nd values of a-Si:D films deposited from pure SiD4 at a pressure of 30 mTorr by variation of the rf power and flow rate. The decrease of Nd at the same Rd in the deuterated samples compared to the hydrogenated ones is caused by the slower thermal desorption rate of D2 compared to H2 at the same temperature. This is suggested by the upward shift of the temperature from which Rd increases in a-Si:D films compared to a-Si:H films, under nominally identical plasma conditions [5], and has recently been observed directly by IRAS as reported by Toyoshima et. al. elsewhere in this volume, The data for a-Si:D can be fitted using eqn. 6b as shown by the lower line in fig. 2. We have simply assumed that CD is less than CH by a factor corresponding to a ~20°C difference which may be ascribed to a difference in attempt frequencies. This demonstrates the imtered systematically by isotopic substitution. Bulk hydrogen diffusion mediated defect densities would imply an increase of Nd in a-Si:D due to slower diffusion of the heavier isotope. The lines in figure 2 tend to saturate at higher Rd. This is due to the increase of abstraction rate of -Sill3 (reaction 1) compared to the thermal desorption rate of hydrogen (reaction 4) so that the system reverts to eqn. 6a from eqn. 6b at higher Rd. The Nd values at high Rd are lower at 400°C than at 250°C because the hopping rate of -Sill3 is thermally enhanced (Cs in eqn. 7b), and Ca/Cs is smaller in eqn. 6a. It would be interesting to verify this trend using a deposition system that overcomes the growth rate limitations of the one we have used. At very low Rd values at Ts=400°C, Ns becomes large and one would expect to see the effect of the DB recombination reaction (5) when Nd would become independent of Rd because then Ns 2 = CHNH2/Cdbr.
(6c)
This can easily be verified experimentally and would yield the activation energy of Cdbr.
On increasing Ts to 450°C, the rate of hydrogen desorption would increase as the value of CH would be expected to increase through the activation energy E H. From eqn. 6b, this increases Nd by the activation energy difference between CH and Cs (~0.9eV). The Nd of sampies prepared at 450°C using procedures similar to those in figs. 1&2 have been plotted in fig. 3. Again, the lines are fits using eqn. 6b and the same values of the parameters indicated above. The Nd values for the deuterated films are lower by the same factor at 450°C as at 400°C. This set of results highlights the fact that the simple surface defect formation process is in remarkable agreement with variation of defect densities in a-Si:H as a function of the growth rate.
.--..
1017~ :
F: 016 ~ 1 "o Z 15 1010-1
~
,, _N-
10 o
_ _,...,,~ m a-~l:N I
10 1
102
Rd ()Vs) Figure 3. The defect density, Nd vs. the growth rate, Rd for a-Si:H & a-Si:D films at Ts=450°C. 3.2. Temperature dependence One of the most consistent trends observed for the variation of deposition parameters is the Ts dependence of Nd which is shown in fig.4 for a-Si:H samples prepared using pure silane at an R d = l k s . According to eqn. 6a, the density of dangling bonds on the surface increases with the abstraction rate constant, Ca, and decreases with the sticking rate constant of the precursors, Cs. Substituting eqns. 7a and 7b in eqn. 6a, Ns/N o = vaexp{-Ea/kT}/vhexp{-Eh)/kT},
(6a')
which says that the fraction of DB sites on the surface (Ns/No) decreases with increasing hop-
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si.'H
ping rate but increases with the rate of abstraction by the precursors. We see from fig.4 that 1 8 , .
10
.,
of Nd with Rd at low Ts has been reported elsewhere in this volume by Kamei et. al. 3.3. Thermal energization of precursors
~
....,_ -/ A/~) C~X\ ,E;]'" ~ (%, ¢~/.v " ~ ~ ° .. : ~ 141.................. ~ ....... ""'"'" ........ . ..... 10 50 150 250 350 450 Ts (°C)
17 ~,-~'10 016 -o 1 Z 15
Figure 4. The defect density, Nd vs. the substrate temperature, Ts. The line, dots and pluses are calculated values (see text), as the temperature increases from, say room temperature, Nd decreases. From eqn. 6a', this implies that Eh>Ea so that Ns/No decreases with increasing Ts. Microscopically, this means that as Ts increases, a -Sill3 on a hydrogenated site on the surface has a larger probability of hopping away than abstracting the hydrogen at that site. The trend of decreasing Nd with increasing Ts calculated using eqn. 6a' is shown by dots in fig. 4. The effect of surface hydrogen desorption (reaction 4) calculated using eqn. 6b is shown by the pluses. The sum of the defect density due to the two reactions is shown by the line in fig. 4. At Ts >350°C, the effect of eqn. 6b dominates, while below 250°C, the effect of eqn. 6a dominates. Therefore, the variation of Nd with Ts over the entire range of Ts is consistent with the defect formation process on the growth surface. At low Ts (<200°C), Nd becomes high and the effect of the DB recombination reaction (5) is significant at Rd values less than ~ l ~ s so that Ns 2 = Ca NH 0 / Cdbr
35
(6d)
whereby, Nd decreases with Rd. The decrease
In fig. 4, we see that Nd is high at both low and high Ts. The PADS method provides one approach to move from the eqn. 6b determined growth regime to the eqn. 6a determined growth regime at high Ts using high Rd (fig.4). Another method that allows reduction of Nd at low Ts is what we call defect reduction by energized precursors (DREP)[6]. DREP separates the temperature responsible for the abstraction reaction from that for precursor hopping in eqn. 6a. Consider the fact that the gas phase temperature decreases away from the substrate in a conventional PECVD reactor [7]. This implies that the precursors produced in the plasma diffuse through an increasing temperature gradient and land on the growing surface. If the precursors react before they can thermalize on the surface, then the temperature of the precursors is less than the growth surface temperature. Although we have assumed that the temperatures in the numerator and denominator of eqn. 6a' are identical, actually, the temperature for hopping is less than that for abstraction. Therefore, by eqn. 6a', the defect density in usual aSi:H films are higher than they might be. We have modified our deposition system by introduction of a mesh type electrode, which can be heated by passing current through it, placed ~2cm above the substrate while the cathode is 4cm above the mesh electrode. Increasing the mesh electrode temperature is expected to change the gas temperature profile above the substrate. The effect of changing the temperature of the mesh electrode (Tin) on Nd, keeping the substrate temperature constant at 150°C, is shown in fig. 5a. At low Tm, the ternperature decreases away from the substrate while at higher Tin, the temperature increases away from the substrate. Consequently, with increasing Tm, the precursors arrive at the substrate with additional thermal energy that enhances their hopping rate. However, the abstraction rate depends primarily on Ts and therefore, with increasing Tin, Ca remains the
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si:H
36
same while Cs increases. Hence by eqn. 6a', Nd decreases with increasing Trn as shown by the line in fig. 5a, which is calculated using the values of parameters given in section 3.1. It is also assumed that the temperature changes linearly from Ts to Tm, the temperature and pressure dependent mean free path (X) of -Sill3 is same as that of silane, and that the effective temperature of -Sill3 is that at 3;L from the substrate, which may be due to multiple collisions being required for gas phase thermalization, Evidence that Ts does not change significantly when Tm increases is seen from the plot of the optical band gap (Eo) versus Tm in fig. 5b. Eo was measured on the same samples used to determine the defect densities plotted in fig. 5a. On the other hand, when Ts increases, Eo decreases, as shown in fig. 5b for samples deposited without the mesh electrode, 17 10
z [ a-Si:H I Ts=150°C
1016
~
F:
""~
I
Thus, we see how small perturbations of the surface reaction processes can lead to orders of magnitude changes of the defect density in a-Si:H. 4. CONCLUSIONS The defect formation model where dangling bonds remaining on the steady state growth surface, by balance of creation and saturation rates, are incorporated into the film bulk as it grows has been used to explain quantitatively the growth rate and substrate temperature dependence of the defect density in a-Si:H. The defect density has been shown to decrease by more than an order of magnitude through an increase of the deposition rate at constant, high substrate temperatures. A reduction of the defect density by two orders of magnitude at a constant, low substrate temperature and constant deposition rate has been attained by subtle control of the surface processes. Furtherrefinementofthemodel, toinclude for example the effects of hydrogen radicals or ions, are possible and will yield greater insight into the growth process of a-Si:H.
A
(a)
REFERENCES J
.
J
~
.,.......... ,..... ! .... : _ . . " [ , . ~ ,. . . 1. .
1.85
,A I II
O
LU 1.75
(b)
Z,
1. N. Itabashi, N. Nishiwaki, M. Magane, S. Naito, T. Goto, A. Matsuda, C. Yamada and E. Hirota, Jpn. J. Appl. Phys. 29
(1990)L505.
• i-
•
1.65 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50 150 250 350 450 T m / T s (°C) Figure 5. (a) The defect density, Nd vs. mesh temperature, Tm, and the line is calculated (see text); (b) the optical gap, Eo vs. mesh temperature Tm ( / k ) for the samples in (a), and Eo vs. substrate temperature, Ts ( A ) for samples prepared without the mesh electrode.
2. Y. Toyoshima, K. Arai, A. Matsuda, and K. Tanaka, J. Non-Cryst. Solids 137/138 (1991) 765. 3. J.-L. Guizot, K. Nomoto, and A. Matsuda, Surf. Sci. 244 (1991) 22. 4. S. Das Sarma, J. Vac. Sci .and Technol. A8 (1990) 2714. 5. G. Ganguly and A. Matsuda, Jpn. J. Appl. Phys. 31 (1992) L1269; Phys. Rev. B47 (1993) 3661. 6. H. Nishio, G. Ganguly, and A. Matsuda, in Proc. Mater. Res. Soc. Spring Meeting 1993, San Francisco, USA (in press). 7. N. Hata, A. Matsuda, and K. Tanaka, Proc. Int. Symp. Plasma Chem. (IUPAC, Tokyo 1987) p-500.
JOU~NALO~ ~ ' ~ d ~ ~LII~
Section 2. Deposition
I m p o r t a n c e of surface p r o c e s s e s in defect formation in a-Si:H Gautam GANGULY and Akihisa MATSUDA Electrotechnical Laboratory, 1-1-4 Umezono, Tsukuba City, Ibaraki, 305 Japan. The deposition temperature and growth rate dependence of the defect density in a-Si:H is shown to depend on the balance between rates of microscopic dangling bond creation and saturation reactions on the steady state growing surface. The surface dangling bonds are incorporated into the bulk as the film grows. On the basis of this model we show that the defect density can be reduced through precursor assisted defect suppression (PADS) at high substrate temperatures (>350°C), and defect reduction by energized precursors (DREP) at low substrate temperatures (<200°C).
1. INTRODUCTION
2. SURFACE REACTIONS
The properties of PECVD a-Si:H films are a sensitive function of the process conditions, Progress with in-situ characterization of the plasma [1] and the growing surface [2] have allowed us considerable insight into the complex growth process. A link between growth conditions and material properties would allow control of material properties. The defect density (Nd) is the key property controlling the functioning of photovoltaic and other devices made using a-Si:H. Nd is empirically found to vary with substrate temperature (Ts) and the deposition rate (Rd) both of which are of critical importance for device fabrication. In this paper we identify the important microscopic reactions, between the precursors from the plasma and the growing surface that create and saturate dangling bonds (DB). We argue that, during steady state growth, the DB creation and saturation rates result in time independent dangling bond densities on the growing surface, and that these DB are incorporated into the bulk as the film grows. This scheme of defect formation on the growing surface explains quantitatively the experimentally observed Rd and Ts dependences of Nd. Replacing hydrogen with deuterium decreases Nd systematically, confirming a sur-
In-situ diode laser absorption spectroscopy reveals that the number density profile and the consequent flux of -Sill3 in a low pressure silane discharge correlates well with the measured deposition rate given the measured sticking probability of 10% [1]. In-situ infrared absorption spectroscopy (IRAS) of the growing aSi:H surface shows that it is saturated with a monolayer of hydrogen at Ts between room temperature and 480°C [2]. Based on these two results we make the simplifying assumption that growth occurs from -Sill3 arriving on a hydrogen covered growth surface. The reactions of -Sill 3 on the growth surface have been speculated upon previously to explain the measured temperature dependence of the surface reaction probability (13)and sticking probability (s) on the a-Si:H growth surface [3]. The surface hydrogen abstraction reaction of -Sill3 creates a surface DB (Si-) and a silane molecule (Sill4)
face mechanism. Two methods of controlling Nd, at high Ts using PADS, and at low Ts using DREP, are presented.
leads to growth by formation of a Si-Si bond. The recombination reaction between -Sill3
- Sill 3 + Si-H = Si- + Sill 4 ,
(1)
while the DB saturation reaction of -Sill3 -SiH3 + Si- = Si-SiH3,
0022-3093/93/$06.00 © 1993 - Elsevier Science Publishers B.V. All rights reserved.
(2)
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si:H
32
-Sill3 + -Sill3 = Si2H6
(3)
forms a disilane molecule (Si2H6) that can desorb from the surface. In addition to these three reactions of -Sill3, there is the thermal hydrogen desorption reaction Si-H + Si-H = 2Si- + H2,
nated site to another until it reacts. We assume also that when a -Sill3 hops on to a DB site, it reacts. The rate constants of reactions 1 and 2 can be written as [3] Ca = B (1/N o) Avaexp{-Ea/kT} , 1-(N H/No) AVhexp{-E h/kT}
(7a)
C s = B (NH/No 2) AVhexp{-E h/kT},
(7b)
(4)
which creates two surface DB and a hydrogen molecule (H2), and the recombination reaction
1-(NH/No) AVhexp{-E h/kT}
between two DB on the surface that diffuse to suitable locations via migration (tunneling!) of surface bonded hydrogen,
where A and B are normalization constants for the total reaction probability of -Sill 3 on a single site and the total reaction probability of -Sill3
Si- + Si- = Si-Si.
during its lifetime on the surface [3]. Eh is the activation energy for hopping of precursors while Ea is the activation energy for the abstraction reaction 1. Vh and Va are the corresponding attempt frequencies. If we assume Arrhenius type processes for reactions 4 & 5 as
(5)
This reaction creates a Si-Si bond on the surface but does not constitute growth. The reactions 1&4 create while reactions 2&5 saturate DB on the growth surface, but the importance of reaction 3 is because it leads to a flux independent precursor diffusion length, unlike other systems [4]. The -Sill3 concentration on the surface during steady state growth, e = [3~ where ~ is the flux of -Sill3, and x is the reaction lifetime of -Sill3 on the growing surface that is determined by the sum of the rates of reactions 1,2 & 3 [3].
3. GROWTH SURFACE TO BULK DEFECTS We now consider the rates of DB creation and saturation due to the surface reactions outlined above. During steady state growth, the rate of change of the surface DB density dNs/dt = CaNH0 - CsNs0 + CHNH 2 - CdbrNs 2 =0
(6)
where the density of hydrogenated surface sites, NH = No, the total density of surface sites, except for Ns, the orders of magnitude smaller density of DB sites. The C's are the temperature dependent rate constants of reactions 1,2,4 & 5. We assume that when it arrives on the growing surface, a -Sill3 that physisorbs, begins to diffuse by hopping from one hydroge-
CH =aH 2 VH exp{-EH/kT}
and
Cdbr = adbr2 Vdbr exp{-Edbr/kT},
(7c) (7d)
where the a's, v's and E's are characteristic distances, attempt frequencies and activation energies, respectively, then we have in eqn.6, a first order approximation to the steady state DB concentration on the growing surface. The connection between the DB density on the growing surface and the film bulk is established through the hypothesis that the dangling bonds on the steady state growing surface are incorporated into the bulk as the film grows. Then, we have the bulk DB density Nd = (Ns)3/2. The macroscopic growth rate Rd = kCsNs0, where k is a conversion factor. Equation 6 can be extended to include the effects of other precursors, Hradical reactions under microcrystalline silicon formation conditions, impurities and ion bombardment. 3.1. Growth rate dependence Equation 6 is the general equation for obtaining the steady state DB density on the growing surface. However, we do not a priori have values of the activation energies of these different processes. We note that when e is
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si.'H
high enough, the first two terms dominate e q u a tion 6. On the other hand, the third term is important when 0 (Rd) is low and Ts is h i g h . Thermal desorption of surface hydrogen has been observed, by IRAS of the growth surface, to occur only at Ts>350°C. The fourth term of eqn. 6 becomes important when 0 (Rd)is low and Ns (Nd) is high. Therefore, at Ts=250°C, we may neglect the third and fourth terms of eqn. 6 to get Ns = NH Ca/Cs.
(6a)
Equation6aimpliesthatNdshouldbeindependent of 0 and hence, Rd. The variation of Nd, evaluated using CPM, as a function of Rd is shown in fig.1 for a-Si:H samples deposited at
Z
f_ ~ . _ ~kLL{ 15r ...~,~.] . . 1010-i
100
increasingRd. The variation of Nd with Rd at 400°C is shown in fig. 2 for a-Si:H samples deposited from pure silane at a pressure of 30 mTorr by changing the flow rate and rf power.
~
1017 . ~.-. . . . . . . . . ~ ~I • a-Si:HJi il , - ~ : 1160 ~ ~ D . O -
E o
.
~ 1015 '10 ,a-Si:Dq 14 ........ 101(~1 100
~
~
101 '
I
I
•
|===
102
Figure 2. The defect density, Nd vs. the growth rate, Rd for a-Si:H & a-Si:D films at Ts=400°C. 101
102
Rd (/VS) Figure 1. The defect density, Nd vs. the g r o w t h rate, Rd for a-Si:H films at Ts=250°C. 250°C under conditions that strive to m a i n t a i n -Sill3 as the major precursors. This is attempted by keeping the pressure at 30mTorr and adjusting the flow rate of pure silane along with the rf power to avoid formation of higher silane related precursors, which is the likely cause of disparate results in previous studies. Above 350°C, the surface bonded hydrogen desorbs thermally as seen by IRAS, and is the reason for the onset of an increase of R d with increasing Ts, keeping plasma conditions constant[3]. Therefore, the effect of reaction 4 becomes significant compared to reaction 1. Hence, neglecting the first term of eqn. 6 cornpared to the third term, we have Ns = NH2 CH / Cs e.
33
(6b)
Thus, at constant Ts, Nd should decrease with
Distinct from the variation at 250°C we see a clear decrease of Nd with increasing Rd. This can be understood as follows: at a constant temperature,the surface hydrogen is thermally evolved and dangling bonds appear on the surface at a constant rate. These surface DB are saturated by the precursors which arrive on the growth surface. Increasing the rate at which the precursors arrive on the surface suppresses the net density of DB remaining on the surface so that we have precursor assisted defect suppression(PADS)[5]. The data in fig. 2 are in agreement with eqn. 6b provided Rd is proportional to 0. Assumingfurther that NH=No=1015cm2, Ea =0.1eV, Eh=0.4eV, Vh/V a ~10 8, aH2VH/VsXB =10-12cm 2, where "~is independent of Ts & Rd, EH=I.3eV from the measured activation energy of Rd [5], the measured value of ~ = 3.2x1015cm-2s-1 at Rd =l,~s [1] along with the average value of N d at 250°C shown by the dashed line in fig.l, we may obtain fits between Nd and Rd using eqn.6b as shown by the upper
34
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si:H
line in fig. 2. Also shown in fig. 2 are the Nd values of a-Si:D films deposited from pure SiD4 at a pressure of 30 mTorr by variation of the rf power and flow rate. The decrease of Nd at the same Rd in the deuterated samples compared to the hydrogenated ones is caused by the slower thermal desorption rate of D2 compared to H2 at the same temperature. This is suggested by the upward shift of the temperature from which Rd increases in a-Si:D films compared to a-Si:H films, under nominally identical plasma conditions [5], and has recently been observed directly by IRAS as reported by Toyoshima et. al. elsewhere in this volume, The data for a-Si:D can be fitted using eqn. 6b as shown by the lower line in fig. 2. We have simply assumed that CD is less than CH by a factor corresponding to a ~20°C difference which may be ascribed to a difference in attempt frequencies. This demonstrates the imtered systematically by isotopic substitution. Bulk hydrogen diffusion mediated defect densities would imply an increase of Nd in a-Si:D due to slower diffusion of the heavier isotope. The lines in figure 2 tend to saturate at higher Rd. This is due to the increase of abstraction rate of -Sill3 (reaction 1) compared to the thermal desorption rate of hydrogen (reaction 4) so that the system reverts to eqn. 6a from eqn. 6b at higher Rd. The Nd values at high Rd are lower at 400°C than at 250°C because the hopping rate of -Sill3 is thermally enhanced (Cs in eqn. 7b), and Ca/Cs is smaller in eqn. 6a. It would be interesting to verify this trend using a deposition system that overcomes the growth rate limitations of the one we have used. At very low Rd values at Ts=400°C, Ns becomes large and one would expect to see the effect of the DB recombination reaction (5) when Nd would become independent of Rd because then Ns 2 = CHNH2/Cdbr.
(6c)
This can easily be verified experimentally and would yield the activation energy of Cdbr.
On increasing Ts to 450°C, the rate of hydrogen desorption would increase as the value of CH would be expected to increase through the activation energy E H. From eqn. 6b, this increases Nd by the activation energy difference between CH and Cs (~0.9eV). The Nd of sampies prepared at 450°C using procedures similar to those in figs. 1&2 have been plotted in fig. 3. Again, the lines are fits using eqn. 6b and the same values of the parameters indicated above. The Nd values for the deuterated films are lower by the same factor at 450°C as at 400°C. This set of results highlights the fact that the simple surface defect formation process is in remarkable agreement with variation of defect densities in a-Si:H as a function of the growth rate.
.--..
1017~ :
F: 016 ~ 1 "o Z 15 1010-1
~
,, _N-
10 o
_ _,...,,~ m a-~l:N I
10 1
102
Rd ()Vs) Figure 3. The defect density, Nd vs. the growth rate, Rd for a-Si:H & a-Si:D films at Ts=450°C. 3.2. Temperature dependence One of the most consistent trends observed for the variation of deposition parameters is the Ts dependence of Nd which is shown in fig.4 for a-Si:H samples prepared using pure silane at an R d = l k s . According to eqn. 6a, the density of dangling bonds on the surface increases with the abstraction rate constant, Ca, and decreases with the sticking rate constant of the precursors, Cs. Substituting eqns. 7a and 7b in eqn. 6a, Ns/N o = vaexp{-Ea/kT}/vhexp{-Eh)/kT},
(6a')
which says that the fraction of DB sites on the surface (Ns/No) decreases with increasing hop-
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si.'H
ping rate but increases with the rate of abstraction by the precursors. We see from fig.4 that 1 8 , .
10
.,
of Nd with Rd at low Ts has been reported elsewhere in this volume by Kamei et. al. 3.3. Thermal energization of precursors
~
....,_ -/ A/~) C~X\ ,E;]'" ~ (%, ¢~/.v " ~ ~ ° .. : ~ 141.................. ~ ....... ""'"'" ........ . ..... 10 50 150 250 350 450 Ts (°C)
17 ~,-~'10 016 -o 1 Z 15
Figure 4. The defect density, Nd vs. the substrate temperature, Ts. The line, dots and pluses are calculated values (see text), as the temperature increases from, say room temperature, Nd decreases. From eqn. 6a', this implies that Eh>Ea so that Ns/No decreases with increasing Ts. Microscopically, this means that as Ts increases, a -Sill3 on a hydrogenated site on the surface has a larger probability of hopping away than abstracting the hydrogen at that site. The trend of decreasing Nd with increasing Ts calculated using eqn. 6a' is shown by dots in fig. 4. The effect of surface hydrogen desorption (reaction 4) calculated using eqn. 6b is shown by the pluses. The sum of the defect density due to the two reactions is shown by the line in fig. 4. At Ts >350°C, the effect of eqn. 6b dominates, while below 250°C, the effect of eqn. 6a dominates. Therefore, the variation of Nd with Ts over the entire range of Ts is consistent with the defect formation process on the growth surface. At low Ts (<200°C), Nd becomes high and the effect of the DB recombination reaction (5) is significant at Rd values less than ~ l ~ s so that Ns 2 = Ca NH 0 / Cdbr
35
(6d)
whereby, Nd decreases with Rd. The decrease
In fig. 4, we see that Nd is high at both low and high Ts. The PADS method provides one approach to move from the eqn. 6b determined growth regime to the eqn. 6a determined growth regime at high Ts using high Rd (fig.4). Another method that allows reduction of Nd at low Ts is what we call defect reduction by energized precursors (DREP)[6]. DREP separates the temperature responsible for the abstraction reaction from that for precursor hopping in eqn. 6a. Consider the fact that the gas phase temperature decreases away from the substrate in a conventional PECVD reactor [7]. This implies that the precursors produced in the plasma diffuse through an increasing temperature gradient and land on the growing surface. If the precursors react before they can thermalize on the surface, then the temperature of the precursors is less than the growth surface temperature. Although we have assumed that the temperatures in the numerator and denominator of eqn. 6a' are identical, actually, the temperature for hopping is less than that for abstraction. Therefore, by eqn. 6a', the defect density in usual aSi:H films are higher than they might be. We have modified our deposition system by introduction of a mesh type electrode, which can be heated by passing current through it, placed ~2cm above the substrate while the cathode is 4cm above the mesh electrode. Increasing the mesh electrode temperature is expected to change the gas temperature profile above the substrate. The effect of changing the temperature of the mesh electrode (Tin) on Nd, keeping the substrate temperature constant at 150°C, is shown in fig. 5a. At low Tm, the ternperature decreases away from the substrate while at higher Tin, the temperature increases away from the substrate. Consequently, with increasing Tm, the precursors arrive at the substrate with additional thermal energy that enhances their hopping rate. However, the abstraction rate depends primarily on Ts and therefore, with increasing Tin, Ca remains the
G. Ganguly, A. Matsuda / Surface processes in defect formation in a-Si:H
36
same while Cs increases. Hence by eqn. 6a', Nd decreases with increasing Trn as shown by the line in fig. 5a, which is calculated using the values of parameters given in section 3.1. It is also assumed that the temperature changes linearly from Ts to Tm, the temperature and pressure dependent mean free path (X) of -Sill3 is same as that of silane, and that the effective temperature of -Sill3 is that at 3;L from the substrate, which may be due to multiple collisions being required for gas phase thermalization, Evidence that Ts does not change significantly when Tm increases is seen from the plot of the optical band gap (Eo) versus Tm in fig. 5b. Eo was measured on the same samples used to determine the defect densities plotted in fig. 5a. On the other hand, when Ts increases, Eo decreases, as shown in fig. 5b for samples deposited without the mesh electrode, 17 10
z [ a-Si:H I Ts=150°C
1016
~
F:
""~
I
Thus, we see how small perturbations of the surface reaction processes can lead to orders of magnitude changes of the defect density in a-Si:H. 4. CONCLUSIONS The defect formation model where dangling bonds remaining on the steady state growth surface, by balance of creation and saturation rates, are incorporated into the film bulk as it grows has been used to explain quantitatively the growth rate and substrate temperature dependence of the defect density in a-Si:H. The defect density has been shown to decrease by more than an order of magnitude through an increase of the deposition rate at constant, high substrate temperatures. A reduction of the defect density by two orders of magnitude at a constant, low substrate temperature and constant deposition rate has been attained by subtle control of the surface processes. Furtherrefinementofthemodel, toinclude for example the effects of hydrogen radicals or ions, are possible and will yield greater insight into the growth process of a-Si:H.
A
(a)
REFERENCES J
.
J
~
.,.......... ,..... ! .... : _ . . " [ , . ~ ,. . . 1. .
1.85
,A I II
O
LU 1.75
(b)
Z,
1. N. Itabashi, N. Nishiwaki, M. Magane, S. Naito, T. Goto, A. Matsuda, C. Yamada and E. Hirota, Jpn. J. Appl. Phys. 29
(1990)L505.
• i-
•
1.65 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50 150 250 350 450 T m / T s (°C) Figure 5. (a) The defect density, Nd vs. mesh temperature, Tm, and the line is calculated (see text); (b) the optical gap, Eo vs. mesh temperature Tm ( / k ) for the samples in (a), and Eo vs. substrate temperature, Ts ( A ) for samples prepared without the mesh electrode.
2. Y. Toyoshima, K. Arai, A. Matsuda, and K. Tanaka, J. Non-Cryst. Solids 137/138 (1991) 765. 3. J.-L. Guizot, K. Nomoto, and A. Matsuda, Surf. Sci. 244 (1991) 22. 4. S. Das Sarma, J. Vac. Sci .and Technol. A8 (1990) 2714. 5. G. Ganguly and A. Matsuda, Jpn. J. Appl. Phys. 31 (1992) L1269; Phys. Rev. B47 (1993) 3661. 6. H. Nishio, G. Ganguly, and A. Matsuda, in Proc. Mater. Res. Soc. Spring Meeting 1993, San Francisco, USA (in press). 7. N. Hata, A. Matsuda, and K. Tanaka, Proc. Int. Symp. Plasma Chem. (IUPAC, Tokyo 1987) p-500.