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Effects of deposition temperature on properties of r.f. glow discharge amorphous silicon thin films
Thin Solid Films. 205 ( 199 I ) I4& I45
140
Effects of deposition temperature amorphous silicon thin films J. L. Andtijar,
E. Bertran, Depurtament (Received
A. Canillas,
de Fisica Aplicadu i Electthica,
January
25. 199 1; accepted
on properties of r.f. glow discharge
C. Roth,
Ihiversitut
J. Serra and G. Sardin
de Barcelonu, Av.Diagonal647,
EO8028 Burceiona (Spain)
May 28, 199 I )
Abstract Thin films of hydrogenated amorphous silicon (a-Si:H) were grown on different substrates by r.f. glow discharge decomposition of silane gas. Hydrogen content and structural properties of films, determined by IR spectroscopy, have been compared with optical properties and density of localized states determined by UV-visible ellipsometry and photothermal deflection measurements respectively. New evidence about the influence of the deposition temperature on the structure and properties of a-Si:H, through its action during the growth process, has been analysed. Some effects induced by deposition temperature on a-Si:H films are the linear decrease in both the hydrogen content and the gap energy as the deposition temperature rises to 400 ‘C as well as the decrease in the density of localized states and the Urbach edge as the temperature increases to 350 “C. Moreover, changes in deposition temperature induce changes in hydrogen incorporation, as has been evidenced through a microstructural parameter calculated from IR absorption peaks at 2000 and 2090 cm-‘. These results suggest that the main mechanism operating in the growth process is thermal activation of the surface mobility which is limited by the hydrogen surface desorption. Above 350 “C, a lower hydrogen coverage arises during the a-Si:H deposition process.
1. Introduction Hydrogenated amorphous silicon (a-Si:H) mainly differs from crystalline silicon in having a high density of localized states at the band edges. For the present, improvements in the quality of a-Si:H have been based on the removal of dangling bonds by hydrogenation. However, in spite of the relatively high hydrogen concentration, which usually is greater than l”/b (about lO*i cmw3), an ubiquitous dangling bonds background remains (around 1017 cmm3). Hydrogen excess affects the SiH, bonds and the density of localized states in the mobility band gap, as well as the electrical and optical properties. Moreover, the hydrogen content has repercussions on the stability of a-Si:H. In spite of that, a-Si:H is being successfully employed as a semiconductor in many applications [I]. The deposition temperature has an explicit influence on the structure and properties of r.f. plasma-deposited a-Si:H films, because it directly takes part in the surface growth process, which constitutes the most determining step in the configuration of the amorphous network. Some effects induced by changes in deposition temperature on a-Si:H films include the following: the decrease in the hydrogen content, in glow discharges, as the deposition temperature rises [2-41, the formation of
polymeric structures (SIN, and SiH,) at low deposition temperatures [5,6], the decrease in the density of localized states with increasing deposition temperature from 300 to 570 K [7,8], the increase in the optical gap energy [2] and in the Urbach edge energy [9] as the temperature decreases, the increase in optical absorption as deposition temperature increases [lo], and some effects on deposition rate [ll]. Apart from the change in the absolute hydrogen content there are also qualitative changes in hydrogen incorporation. From thermal desorption spectroscopy a two-peak spectrum can be obtained for low deposition temperature through the hydrogen effusion [12], where the low temperature hydrogen effusion peak has been ascribed to the presence of polysilane-like void-rich material [13, 141. This may be a source of the Staebler Wronski effect [15]. In this paper, the hydrogen content and structural properties of a-Si:H thin films are compared with optical properties and density of localized states. New evidence about the influence of the deposition temperature on the structure and properties of a-Si:H thin films, through its action during the growth process, is reported. Hydrogen incorporation is discussed in terms of a thermal activation of the mobility of adsorbed radical species on the growing surface.
E. Bertran et al. / R.f glow discharge a-S
2. Experimental features 2.1. Deposition process Thin films of a-Si:H were prepared by r.f. glow discharge decomposition of pure silane gas. The deposition reactor [ 161consists of two vertical electrodes with 40 mm gap capacitively coupled to an r.f. power supply working at 13.56 MHz. An automatic matching network minimizes reflected power coming from the load impedance of the discharge. The output r.f. power was maintained constant at 8 W. This power corresponds to a cathode power density of 20 mW cmm2. Films of a-Si:H were deposited on several substrates such as crystalline silicon (c-Si) wafers, quartz, Corning 7059 glass and NiCr, in order to perform different characterizations. The deposition temperature was varied from 200 “C to 400 “C for samples grown at 30 Pa of pressure, and from 100 “C to 350 “C for samples grown at 3 Pa. The silane flow rate was 30 standard cm3 min- ‘. The silane pressure was measured by an absolute pressure capacitive meter and it was kept nearly constant during the deposition process. The substrates were protected by a shutter during plasma stabilization. 2.2. Ellipsometer The reactor chamber is equipped with a fast phasemodulated ellipsometer which can operate in the UV-visible range in two ways: the real-time mode and the spectroscopic mode [I 71. Ellipsometry measures the angles (Y,d) as a function of wavelength. These angles (Y,d) are related to the complex reflectance ratio p = ?Jfs between the parallel (p) and perpendicular (s) reflection coefficients of the electric field by p”= tan Y eiA
(1)
In the case of a homogeneous material with a sharp interface with the ambient, the complex pseudo-dielectric function E”is related to 6 through the relationship: E’= .sl -iE2 = sin2q [l +{(l -p’)/(l +i?)}2 tan2q]
(2)
where cp is the incidence angle of the light beam on the sample surface. The real-time mode provides the evolution of the complex reflectance ratio during the growth process of the films. This permits the growth kinetics of the a-Si:H films to be monitored at a given wavelength. From this evolution, the deposition rate and the film microstructure can be determined by fitting the experimental trajectories (Y/J) to the theoretical growth models [18]. The spectroscopic mode of the ellipsometer reveals the dependence of the complex pseudo-dielectric function 1 on the photon energy. The maximum stmax of the imaginary part of g is related to the material density as well as to the surface roughness and to the oxidation [19].
141
An increase in the maximum value of elmax indicates a qualitative increase in the material density. On the contrary, the s2 spectrum allows us to determine the optical gap energy Eg corresponding to interband transitions between the localized states and the extended states. The EB values have been calculated from the expression derived by Taut: E&21’2 = B(E-E,)
(3)
where E represents the radiation energy and B is a nondimensional constant proportional to the probability of an interband transition [20]. 2.3. IR spectroscopy Transmittance measurements were performed on a-Si:H films with a Fourier transform IR (FTIR) spectrometer (Nicolet SZDX), with the purpose of determining the hydrogen content and the microstructure of the films. The spectra of a-Si:H films were obtained in the range 400-4000 cm-‘, with a resolution of 4 cm-‘, after removing the absorption owing to the c-Si substrate. The absorption coefficient c1can be calculated from the transmittance spectrum [5]. From the wagging absorption mode at 630 cm- ’ the degree of bonded hydrogenation in a-Si:H films was determined [21]; furthermore, from integration of the absorption peaks at 2000 cm- 1 and 2090 cm-‘, the film’s microstructure can be estimated. The microstructure factor R is defined by the ratio [22] R = [2090]/([2000] + [2090])
(4)
where [2000] and [2090] represent the integrated absorption centred at 2000 cm- ’ and 2090 cm- ’ stretching modes respectively. The mode centred at 2000 cm-’ corresponds to individual Si-H bonds, whereas the 2090 cm- ’ peak has been ascribed both to Si-H, bonds and to Si-H bonds assembled on internal surfaces inside microcavities of a-Si:H [21]. In any case, R is proportional to the hydrogen fraction bonded in some kind of microstructure. 2.4. Photothermal deflection spectroscopy Measurements of photothermal deflection spectroscopy (PDS) were performed in order to determine the Urbach edge energy and the evolution of the density of localized states of a-Si:H films with the deposition temperature. This technique measures the calorific energy absorbed by the a-Si:H film-substrate structure on which a monochromatic beam is focused. For this analysis quartz substrates were used because they exhibit a very low optical absorption in the photon energy range considered (0.9-2.0 eV). This technique allows us to calculate weak absorption coefficients of thin films through measurements of the total energy absorbed by the sample. The experimental arrangement used in the present
E. Bertrun et ul. :I R.fI glow discharge a-Si
142
study consists of a light, proceeding from a 40 W halogen lamp, passing through a chopper and a monochromator with 2 nm of resolution, and being focused on the sample. The He-Ne laser beam of power 5 mW, which passes parallel to the sample surface, deflects because of a refractive index gradient of the medium (Ccl,) adjacent to the sample, induced by the surface temperature of the film. This deflection is measured using a pair of light detectors and a lock-in amplifier (Stanford SR530). In order to calibrate the assembly the signal was previously normalized to a standard graphite sample. The normalized PDS signal S for thermally thin samples is given by [23] S = S,(l -emad)
(5)
where c(is the absorption coefficient, d is the film thickness and S, is the PDS signal corresponding to a high photon energy, just where a-Si:H samples exhibit high cc values. Accurate PDS measurements provide the weak absorption coefficient values corresponding to subgap states. The excess absorption rex due to subgap states can be calculated from: a ex = 2 - aOeEIEo
(6)
where E is the photon energy, E, is the Urbach edge energy and a0 a constant. The density of localized states has been calculated from the integration of the optical absorption excess of the band gap [24].
3. Experimental
results and discussion
Some effects on optical and structural properties of the a-Si:H thin films induced by the deposition temperature T, have been analysed. Figure 1 shows the maximum values of E* as a function of the deposition temperature. These
values were obtained by in situ ellipsometric measurements at room temperature, and they correspond to a-Si:H films deposited on NiCr substrates at 8 W of r.f. power. 30 standard cm3 min ’ of silane flow and 30 Pa of pressure. The maximum values eZbmaxof s2 corresponding to bulk material were determined using a multilayer model according to a structure constituted by a homogeneous a-Si:H film with surface roughness [25]. The equivalent thickness of the surface roughness, obtained by kinetic ellipsometry, is around 1.5 nm in all cases. The cZbmax values have been plotted in Fig. I together with the experimental sZmax values showing the effect of surface roughness in ellipsometric measurements. The high values of .sZbmaxobtained for T, > 250 “C indicate that the a-Si:H films possess a very low porosity. in The behaviour of E~,,,=~ us. the deposition temperature Fig. 1 indicates that the films become denser as the deposition temperature increases to 350 “C. At low pressure conditions (3 Pa) (Fig. 2). an increase in 3.7 3.6 '*3 5 3.4 3.3 50
100 150 Deposition
200 250 300 350 Deposition Temperature
400 (“C)
4
4-50
Fig. I. Dependence on deposition temperature of the maximum value of the imaginary part of the dielectric function, determined from spectral ellipsometric measurements. These values correspond to thin films of a-Si:H deposited on NiCr at 30 Pa of pressure. standard cm3 min- ’ of SiH, flow and 8 W of r.f. power. 0, experimental values corresponding to a semi-infinite medium with a surface roughness; 0, a-Si:H bulk material.
350 (“C)
4 0
Fig. 3. Refractive index of a-Si:H for near-IR long wavelengths, determined from transmittance measurements. as a function of deposition temperature of a-Si:H thin films grown on glass substrates at Wofr.f,power. 3 Pa ofpressure. 30 standard cm3 min ’ ofSiH,flowand8
refractive temperature n30 values spectra of Sellmeier’s
index n, for long wavelength with deposition (from 100 to 200 ‘C) has been observed. The were calculated from near-IR transmittance a-Si:H films on glass substrates by means of one-oscillator dispersion relation [26]:
n2 = I+(n,’
161 150
200 250 300 Temperature
- l)/12/(i.2 - ;““I)
(7)
The evolution of n, with deposition temperature, for films grown under low pressure conditions, corroborates the previous interpretation about the density of the a-Si:H films. The material density increase can be explained in terms of the a-Si:H growth model suggested by Tanaka and Matsuda [27]. This model assumes SiH, radicals to be precursor species of a-Si:H deposition. The surface diffusion coefficient of adsorbed SiH, radicals is determined by the deposition temperature and the surface hydrogen coverage factor. Therefore, the observed densification, as the deposition temperature increases to 250 “C, is explained by an increase in surface radical diffusion which favours the formation of a higher compactness of the a-Si:H network. However, over a specific temperature (350 “C) the subsequent increase in
E. Bertran et al. / R.f. glow discharge a-S
temperature eliminates hydrogen from the surface, therefore decreasing the hydrogen coverage factor, which explains the decrease in sZmax at 400 “C in Fig. 1. Figure 3 shows the dependence of the microstructure factor R on the deposition temperature for samples grown at 30 Pa. The microstructure factor reaches a minimum at 350 “C at the same time as Q,,,._ exhibits a maximum at this temperature (Fig. 1). However, the hydrogen content for these samples (Fig. 4), determined by FTIR spectroscopy, decreases linearly from 11% to 5% when the deposition temperature rises from 200 “C to 400 “C without displaying any particular change at 350 “C. In Fig. 5, EB also decreases linearly from 2.02 eV to 1.88 eV as T, increases from 200 “C to 400 “C. The correlation between the hydrogen content and the optical gap energy E, is determined by comparing both Fig. 4 and Fig. 5. It is
0.3
143
extensively admitted that the optical gap energy is affected by hydrogen content [28], and the experimental data from the literature reveal a clear correlation between the optical gap energy and the bonded hydrogen density [29]. Von Roedern et al. [30] concluded, from photoemission experiments, that hydrogen incorporation removes states on the top of the valence band. However, Cody et al. [31] consider that the optical gap energy is determined by the degree of disorder in the material and that hydrogen indirectly affects the gap energy, releasing stress to the amorphous network. In addition to the saturation of dangling bonds, another important effect of hydrogen incorporation arises from its influence on the network cross-linking by lowering its self-coordination (Si-Si bonding). In effect, hydrogen introduces a fourth hetero coordination @i-H), which acts as a bonding terminal, thus leading to a looser network. The variations in the slope B and the energy &, of the Urbach edge with the deposition temperature are shown in Fig. 6 and Fig. 7 respectively. Both dependences
w 0.2 0.1 i
O.O ‘0 150
200 250 300 350 Deposition Temperature
400 (“C)
Fig. 3. Deposition temperature dependence of the microstructure parameter R, defined from the a-Si:H stretching modes. Deposition conditions are 30 Pa of pressure. 30 standard cm3 min-’ of SiH, flow and 8 W of r.f. power.
9
I
150
,
200 250 300 350 Deposition Temperature
400 (“C)
450
Fig. 6. Deposition temperature dependence of the slope B referred to the linear dependence of Ee,‘!2 on energy E of the Taut model. The deposition conditions were 30 Pa of pressure, 30 standard cm3 min- ’ of SiH, flow and 8 W of r.f. power.
k I”
41d 150
200 250 300 350 Deposition Temperature
400 (“C)
450
Fig. 4. Variation with the deposition temperature in the bonded hydrogen content in a-Si:H films, determined from the wagging mode of FTIR spectra. The deposition conditions are 30 Pa of pressure, 30 standard cm3 min _ ’ of SiH, flow and 8 W of r.f. power.
55
1
150
I
I
I
200 250 300 350 Deposition Temperature
I
400 (“C)
450
Fig. 7. Deposition temperature dependence of the Urbach edge energy, determined by PDS of a-Si:H thin films grown on quartz substrates at 30 Pa of pressure, 30 standard cm3 min-’ of SiH, flow and 8 W of r.f. power.
1.6 &I 150
200 250 300 350 Deposition Temperature
400 (“C)
450
Fig. 5. Optical gap energy dependence on deposition temperature of bulk a-Si:H. These values were calculated from dielectric function spectra, determined by elhpsometric measurements, taking into account the Taut model. The deposition conditions are 30 Pa of pressure, 30 standard cm3 min-’ of SiH, flow and 8 W of r.f. power.
indicate that the disorder in the amorphous network reaches a minimum around 300-350 “C and, furthermore, that disorder is higher at low deposition temperatures. The localized states density N, in the gap (Fig. 8), measured by PDS, also exhibits a minimum value in the interval 300-350 “C (N, = 3.7 x 1Or6 cmm3). However,
E. Bertrun et ul. / R.f: glow discharge u-Si
144
150
200 250 300 350 Deposition Temperature
400 (“C)
Fig. 8. Deposition temperature dependence of the density of localized states in the forbidden band of a-Si:H, determined by PDS measurements, The films of a-Si:H were deposited at 30 Pa of pressure. 30 standard cm3 min _ ’ of SiH, flow and 8 W of r.f. power.
the film deposited at 400 “C is the film that exhibits a higher density of states in the gap. The real-time ellipsometric trajectories corresponding to film growth have been analysed assuming a previously suggested model [ 181, where hemispherical nucleation, the coalescence phase and the homogeneous growth are taken into account. Figure 9 shows the film thickness at
:kGT r
Deposition
--330
300
temperature
-4;10 4504 (“C)
Fig, 9. Deposttion temperature dependence of the a-Si:H film thickness at the end of the nucleation and coalescence phases, determined from real-time ellipsometric measurements. The films were deposited at 30 Pa of pressure. 30 standard cm3 min 1of SiH, flow and 8 W of r.f. power.
the end of the coalescence phase as a function of deposition temperature. This parameter decreases to a minimum at a deposition temperature of 350 “C. This behaviour can be a consequence of two phenomena related to hydrogen content and surface hydrogen coverage. The nucleation phase depends on the state of the surface and of the hydrogen coverage, and therefore a delay in the final stage of the nucleation phase as a consequence of the decrease in hydrogen coverage above 350 “C can be expected. On the contrary, the coalescence phase can be completed earlier because of the increase in the dangling bonds between growing cylinders [ 181as the deposition temperature rises. Results in this paper reveal the influence of deposition temperature on different properties of a-Si:H. However, it is possible that this technological parameter does not directly affect the a-Si:H properties, but rather it acts on hydrogen through the deposition process mechanisms. At a deposition temperature of 350 “C, all studied properties exhibit an extreme behaviour. In spite of that, the characterization techniques used are independent of
each other, and the measured properties are based on different principles. This suggests that hydrogen incorporated in the network controls the a-Si:H properties. The important role of hydrogen in the structure of a-Si:H is well known. Hydrogen incorporation during the growth of a-Si:H thin films is controlled by two antagonistic kinetics. The first kinetics is associated with the flux of the impinging hydride radicals. The kinetics of the hydrogen contribution depends on the flux of the impinging radicals and on the number of hydrogen atoms associated with each type. The second kinetics which derives from two heterogeneous reactions is associated with the ejection of hydrogen from hydrogen-rich radicals (hydrogen desorption). These two kinetics compete severely, the resultant leading to a relatively small incorporation of hydrogen (about IO’:;) compared with the H:Si ratio of the impinging species (H:Si 2 300’;$ thus about 290”,, of the hydrogen has been ejected. Since the temperature is a determinant of the heterogeneous reaction. it allows us to “modulate” the hydrogen incorporation. Hydrogen incorporates in amorphous silicon under different bond configurations such as Si-H, Si-HZ. Si-H, and (Si-H,),. depending on preparation conditions. In the light of the results, everything seems to indicate a thermal activation of the surface mobility of adsorbed radical species on the growing film, limited by the desorption of hydrogen above 350 “C. The arguments supporting this assumption are stated as follows. (a) The material density (as a synonym of F~,,,_ in Fig. 1) increases as the deposition temperature increases before reaching the value of 350 ‘C. which suggests that the increase in surface mobility favours the formation of Si-H bonds with respect to other bonding‘forms. (b) The evolution of hydrogen content (Fig. 4) and the decrease in microstructure factor R which means a decrease in Si-H, and 2(Si-H) bonds with respect to the formation of Si-H bond (Fig. 3) are also indicative that films become denser. (c) The Urbach edge energy evolution in Fig. 7, which indicates a reduction in localized states in the band tails as the temperature rises to 350 ‘C, can be interpreted as a higher efficiency of the hydrogen in saturating dangling bonds. This interpretation is supported by the decrease in density of localized states as the temperature rises (Fig. 8). (d) The real-time elliljsometric results also indicate an increase in surface mobility, since both the nucleation phase and the coalescence phase become faster as the deposition temperature rises to 350 C.
4. Conclusion It becomes clear that the deposition temperature is a critical parameter which has beneficial effects up to
E. Bertran et al. / R.f. glow discharge a-Si
350 “C by lowering the microstructure parameter, the material density, the density of localized states and the Urbach edge as the deposition temperature increases. These dependences indicate an increase in surface mobility which favours the formation of Si-H bonds with respect to other bonding forms and a higher efficiency of the hydrogen in saturating dangling bonds. Above this threshold temperature, at which the desorption of hydrogen incorporated in the network becomes important, the density of states, the Urbach edge, the material density and the microstructure parameter increase. The main mechanism active in the growth process is the thermal activation of the surface mobility of adsorbed radical species on the growing surface, which are limited by the hydrogen surface desorption. Above a deposition temperature of 350 “C, less hydrogen coverage occurs during the a-Si:H deposition process.
6 7 8 9 10 11 12 13 14 15
16 17
Acknowledgments
18
The authors acknowledge Professor J. L. Morenza and Dr. J. Andreu for their suggestions and they thank Dr. A. Lloret for fruitful discussions. This work has been supported by the CAICYT de1 Ministerio de Education y Ciencia of Spain under Contract 798/84 and the CICYT of Spain under Contract MAT 955190.
19 20
21 22 23
References P. G. LeComber, J. Non-Crysr. Solids, 115 (1989) 1. .I. Perrin, I. Solomon, B. Bourdon. J. Fontenille and E. Ligeon, Thin Solid Films, 62 (1979) 327. M. Milleville, W. Fuhs, F. J. Demond, H. Mannsperger, G. Miiller and S. Kalbitzer, Appl. Phys. Left., 34 (1979) 173. C. R. Ross and J. Jaklik, J. Appl. Phys., 55 (1984) 3785. M. H. Brodsky, M. Cardona and J. J. Cuomo, Phys. Rev. B. 16 (1977) 3556.
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G. Lucovski, R. J. Nemanich and J. C. Knights, Phys. Rev. B, I9 (1979) 2064. W. E. Spear and P. G. LeComber, J. Non-Cryst. Solids, 8-10 (1972) 727. A. Madan, P. G. LeComber and W. E. Spear, J. Non-Cryst. Solids, 20 (1976) 239. S. Kawasaki, K. Sate, K. Suzuki, H. Takeuchi, K. Kuroiwa and Y. Tarui, Jpn. J. Appl. Phys., 26 (1987) 1400. P. J. Zanzucchi, C. R. Wronski and D. E. Carlson, J. Appl. Phys., 48 (1977) 5227. J. L. Andtijar, E. Bertran, A. Canillas, J. Campmany and J. L. Morenza, J. Appl. Phys., 69 (1991) 3757. G. Sardin, J. Andreu, J. C. Delgado and J. L. Morenza, Sol. Energy Muter., I7 (1988) 227. M. H. Brodsky, M. A. Frisch and J. F. Ziegler, Appl. Phys. Left., 30 (1977) 561. W. Beyer and H. Wagner, J. Phys. Colloq., 42 (1981) 783. G. Sardin, J. Kasaneva, E. Bertran, J. L. Andtijar, C. Roth, J. Andreu and J. L. Morenza, Proc. 9th European Communities Photovoltaic Solar Energy Conf , Freiburg. 1989, Kluwer, Dordrecht, 1989, p. 1021. J. L. Andujar, E. Bertran, A. Canillas, J. Esteve, J. Andreu and J. L. Morenza, Vacuum, 39 (1989) 795. A. Canillas, E. Bertran, J. L. Andujar and J. L. Morenza, Vacuum, 39 (1989) 785. A. Canillas, E. Bertran, J. L. Anditjar and B. Drevillon, J. Appl. Phys., 68 (1990) 2752. B. Drevillon and F. Vaillant, Thin Solid Films, 124 (1985) 217. L. Ley, in J. D. Joannopoulos and G. Lucovsky (eds.), The Physics of Hydrogenated Amorphous Silicon II, Springer, Berlin, 1984, p. 61. H. Shanks, C. J. Fang, L. Ley, M. Cardona, F. J. Demond and S. Kalbitzer, Phys. Status Solidi I?, 100 (1980) 43. A. H. Mahan, P. Raboisson and R. Tsu, Appl. Phys. Lett., 50 (1987) 335. W. B. Jackson, N. M. Amer, A. C. Boccara and D. Fournier, Appl. Opr., 20 (1981) 1333. W. B. Jackson and N. M. Amer, Phys. Rev. B, 25 (1982) 5559. D. E. Aspnes. Thin Solid Films, 89 (1982) 249. F. A. Jenkins and H. E. White, Fundamentals of Optics, McGrawHill, New York, 1957. K. Tanaka and A. Matsuda, Mafer. Sci. Rep., 2 (1987) 139. C. C. Tsai and H. Fritzsche, Sol. Energy Marer., I (1979) 29. R. V. Kruzelecky, D. Rakansky, S. Zukotynski and J. M. Perz, J. Non-Cryst. Solids, 99 (1988) 89. B. von Roedern, L. Ley and M. Cardona, Phys. Rev. Lert.. 39 (1977) 1576. G. D. Cody, T. Tiedje, B. Abeles, B. Brooks and Y. Goldstein, Phys. Rev. Le/t., 47 (1981) 1480.
140
Effects of deposition temperature amorphous silicon thin films J. L. Andtijar,
E. Bertran, Depurtament (Received
A. Canillas,
de Fisica Aplicadu i Electthica,
January
25. 199 1; accepted
on properties of r.f. glow discharge
C. Roth,
Ihiversitut
J. Serra and G. Sardin
de Barcelonu, Av.Diagonal647,
EO8028 Burceiona (Spain)
May 28, 199 I )
Abstract Thin films of hydrogenated amorphous silicon (a-Si:H) were grown on different substrates by r.f. glow discharge decomposition of silane gas. Hydrogen content and structural properties of films, determined by IR spectroscopy, have been compared with optical properties and density of localized states determined by UV-visible ellipsometry and photothermal deflection measurements respectively. New evidence about the influence of the deposition temperature on the structure and properties of a-Si:H, through its action during the growth process, has been analysed. Some effects induced by deposition temperature on a-Si:H films are the linear decrease in both the hydrogen content and the gap energy as the deposition temperature rises to 400 ‘C as well as the decrease in the density of localized states and the Urbach edge as the temperature increases to 350 “C. Moreover, changes in deposition temperature induce changes in hydrogen incorporation, as has been evidenced through a microstructural parameter calculated from IR absorption peaks at 2000 and 2090 cm-‘. These results suggest that the main mechanism operating in the growth process is thermal activation of the surface mobility which is limited by the hydrogen surface desorption. Above 350 “C, a lower hydrogen coverage arises during the a-Si:H deposition process.
1. Introduction Hydrogenated amorphous silicon (a-Si:H) mainly differs from crystalline silicon in having a high density of localized states at the band edges. For the present, improvements in the quality of a-Si:H have been based on the removal of dangling bonds by hydrogenation. However, in spite of the relatively high hydrogen concentration, which usually is greater than l”/b (about lO*i cmw3), an ubiquitous dangling bonds background remains (around 1017 cmm3). Hydrogen excess affects the SiH, bonds and the density of localized states in the mobility band gap, as well as the electrical and optical properties. Moreover, the hydrogen content has repercussions on the stability of a-Si:H. In spite of that, a-Si:H is being successfully employed as a semiconductor in many applications [I]. The deposition temperature has an explicit influence on the structure and properties of r.f. plasma-deposited a-Si:H films, because it directly takes part in the surface growth process, which constitutes the most determining step in the configuration of the amorphous network. Some effects induced by changes in deposition temperature on a-Si:H films include the following: the decrease in the hydrogen content, in glow discharges, as the deposition temperature rises [2-41, the formation of
polymeric structures (SIN, and SiH,) at low deposition temperatures [5,6], the decrease in the density of localized states with increasing deposition temperature from 300 to 570 K [7,8], the increase in the optical gap energy [2] and in the Urbach edge energy [9] as the temperature decreases, the increase in optical absorption as deposition temperature increases [lo], and some effects on deposition rate [ll]. Apart from the change in the absolute hydrogen content there are also qualitative changes in hydrogen incorporation. From thermal desorption spectroscopy a two-peak spectrum can be obtained for low deposition temperature through the hydrogen effusion [12], where the low temperature hydrogen effusion peak has been ascribed to the presence of polysilane-like void-rich material [13, 141. This may be a source of the Staebler Wronski effect [15]. In this paper, the hydrogen content and structural properties of a-Si:H thin films are compared with optical properties and density of localized states. New evidence about the influence of the deposition temperature on the structure and properties of a-Si:H thin films, through its action during the growth process, is reported. Hydrogen incorporation is discussed in terms of a thermal activation of the mobility of adsorbed radical species on the growing surface.
E. Bertran et al. / R.f glow discharge a-S
2. Experimental features 2.1. Deposition process Thin films of a-Si:H were prepared by r.f. glow discharge decomposition of pure silane gas. The deposition reactor [ 161consists of two vertical electrodes with 40 mm gap capacitively coupled to an r.f. power supply working at 13.56 MHz. An automatic matching network minimizes reflected power coming from the load impedance of the discharge. The output r.f. power was maintained constant at 8 W. This power corresponds to a cathode power density of 20 mW cmm2. Films of a-Si:H were deposited on several substrates such as crystalline silicon (c-Si) wafers, quartz, Corning 7059 glass and NiCr, in order to perform different characterizations. The deposition temperature was varied from 200 “C to 400 “C for samples grown at 30 Pa of pressure, and from 100 “C to 350 “C for samples grown at 3 Pa. The silane flow rate was 30 standard cm3 min- ‘. The silane pressure was measured by an absolute pressure capacitive meter and it was kept nearly constant during the deposition process. The substrates were protected by a shutter during plasma stabilization. 2.2. Ellipsometer The reactor chamber is equipped with a fast phasemodulated ellipsometer which can operate in the UV-visible range in two ways: the real-time mode and the spectroscopic mode [I 71. Ellipsometry measures the angles (Y,d) as a function of wavelength. These angles (Y,d) are related to the complex reflectance ratio p = ?Jfs between the parallel (p) and perpendicular (s) reflection coefficients of the electric field by p”= tan Y eiA
(1)
In the case of a homogeneous material with a sharp interface with the ambient, the complex pseudo-dielectric function E”is related to 6 through the relationship: E’= .sl -iE2 = sin2q [l +{(l -p’)/(l +i?)}2 tan2q]
(2)
where cp is the incidence angle of the light beam on the sample surface. The real-time mode provides the evolution of the complex reflectance ratio during the growth process of the films. This permits the growth kinetics of the a-Si:H films to be monitored at a given wavelength. From this evolution, the deposition rate and the film microstructure can be determined by fitting the experimental trajectories (Y/J) to the theoretical growth models [18]. The spectroscopic mode of the ellipsometer reveals the dependence of the complex pseudo-dielectric function 1 on the photon energy. The maximum stmax of the imaginary part of g is related to the material density as well as to the surface roughness and to the oxidation [19].
141
An increase in the maximum value of elmax indicates a qualitative increase in the material density. On the contrary, the s2 spectrum allows us to determine the optical gap energy Eg corresponding to interband transitions between the localized states and the extended states. The EB values have been calculated from the expression derived by Taut: E&21’2 = B(E-E,)
(3)
where E represents the radiation energy and B is a nondimensional constant proportional to the probability of an interband transition [20]. 2.3. IR spectroscopy Transmittance measurements were performed on a-Si:H films with a Fourier transform IR (FTIR) spectrometer (Nicolet SZDX), with the purpose of determining the hydrogen content and the microstructure of the films. The spectra of a-Si:H films were obtained in the range 400-4000 cm-‘, with a resolution of 4 cm-‘, after removing the absorption owing to the c-Si substrate. The absorption coefficient c1can be calculated from the transmittance spectrum [5]. From the wagging absorption mode at 630 cm- ’ the degree of bonded hydrogenation in a-Si:H films was determined [21]; furthermore, from integration of the absorption peaks at 2000 cm- 1 and 2090 cm-‘, the film’s microstructure can be estimated. The microstructure factor R is defined by the ratio [22] R = [2090]/([2000] + [2090])
(4)
where [2000] and [2090] represent the integrated absorption centred at 2000 cm- ’ and 2090 cm- ’ stretching modes respectively. The mode centred at 2000 cm-’ corresponds to individual Si-H bonds, whereas the 2090 cm- ’ peak has been ascribed both to Si-H, bonds and to Si-H bonds assembled on internal surfaces inside microcavities of a-Si:H [21]. In any case, R is proportional to the hydrogen fraction bonded in some kind of microstructure. 2.4. Photothermal deflection spectroscopy Measurements of photothermal deflection spectroscopy (PDS) were performed in order to determine the Urbach edge energy and the evolution of the density of localized states of a-Si:H films with the deposition temperature. This technique measures the calorific energy absorbed by the a-Si:H film-substrate structure on which a monochromatic beam is focused. For this analysis quartz substrates were used because they exhibit a very low optical absorption in the photon energy range considered (0.9-2.0 eV). This technique allows us to calculate weak absorption coefficients of thin films through measurements of the total energy absorbed by the sample. The experimental arrangement used in the present
E. Bertrun et ul. :I R.fI glow discharge a-Si
142
study consists of a light, proceeding from a 40 W halogen lamp, passing through a chopper and a monochromator with 2 nm of resolution, and being focused on the sample. The He-Ne laser beam of power 5 mW, which passes parallel to the sample surface, deflects because of a refractive index gradient of the medium (Ccl,) adjacent to the sample, induced by the surface temperature of the film. This deflection is measured using a pair of light detectors and a lock-in amplifier (Stanford SR530). In order to calibrate the assembly the signal was previously normalized to a standard graphite sample. The normalized PDS signal S for thermally thin samples is given by [23] S = S,(l -emad)
(5)
where c(is the absorption coefficient, d is the film thickness and S, is the PDS signal corresponding to a high photon energy, just where a-Si:H samples exhibit high cc values. Accurate PDS measurements provide the weak absorption coefficient values corresponding to subgap states. The excess absorption rex due to subgap states can be calculated from: a ex = 2 - aOeEIEo
(6)
where E is the photon energy, E, is the Urbach edge energy and a0 a constant. The density of localized states has been calculated from the integration of the optical absorption excess of the band gap [24].
3. Experimental
results and discussion
Some effects on optical and structural properties of the a-Si:H thin films induced by the deposition temperature T, have been analysed. Figure 1 shows the maximum values of E* as a function of the deposition temperature. These
values were obtained by in situ ellipsometric measurements at room temperature, and they correspond to a-Si:H films deposited on NiCr substrates at 8 W of r.f. power. 30 standard cm3 min ’ of silane flow and 30 Pa of pressure. The maximum values eZbmaxof s2 corresponding to bulk material were determined using a multilayer model according to a structure constituted by a homogeneous a-Si:H film with surface roughness [25]. The equivalent thickness of the surface roughness, obtained by kinetic ellipsometry, is around 1.5 nm in all cases. The cZbmax values have been plotted in Fig. I together with the experimental sZmax values showing the effect of surface roughness in ellipsometric measurements. The high values of .sZbmaxobtained for T, > 250 “C indicate that the a-Si:H films possess a very low porosity. in The behaviour of E~,,,=~ us. the deposition temperature Fig. 1 indicates that the films become denser as the deposition temperature increases to 350 “C. At low pressure conditions (3 Pa) (Fig. 2). an increase in 3.7 3.6 '*3 5 3.4 3.3 50
100 150 Deposition
200 250 300 350 Deposition Temperature
400 (“C)
4
4-50
Fig. I. Dependence on deposition temperature of the maximum value of the imaginary part of the dielectric function, determined from spectral ellipsometric measurements. These values correspond to thin films of a-Si:H deposited on NiCr at 30 Pa of pressure. standard cm3 min- ’ of SiH, flow and 8 W of r.f. power. 0, experimental values corresponding to a semi-infinite medium with a surface roughness; 0, a-Si:H bulk material.
350 (“C)
4 0
Fig. 3. Refractive index of a-Si:H for near-IR long wavelengths, determined from transmittance measurements. as a function of deposition temperature of a-Si:H thin films grown on glass substrates at Wofr.f,power. 3 Pa ofpressure. 30 standard cm3 min ’ ofSiH,flowand8
refractive temperature n30 values spectra of Sellmeier’s
index n, for long wavelength with deposition (from 100 to 200 ‘C) has been observed. The were calculated from near-IR transmittance a-Si:H films on glass substrates by means of one-oscillator dispersion relation [26]:
n2 = I+(n,’
161 150
200 250 300 Temperature
- l)/12/(i.2 - ;““I)
(7)
The evolution of n, with deposition temperature, for films grown under low pressure conditions, corroborates the previous interpretation about the density of the a-Si:H films. The material density increase can be explained in terms of the a-Si:H growth model suggested by Tanaka and Matsuda [27]. This model assumes SiH, radicals to be precursor species of a-Si:H deposition. The surface diffusion coefficient of adsorbed SiH, radicals is determined by the deposition temperature and the surface hydrogen coverage factor. Therefore, the observed densification, as the deposition temperature increases to 250 “C, is explained by an increase in surface radical diffusion which favours the formation of a higher compactness of the a-Si:H network. However, over a specific temperature (350 “C) the subsequent increase in
E. Bertran et al. / R.f. glow discharge a-S
temperature eliminates hydrogen from the surface, therefore decreasing the hydrogen coverage factor, which explains the decrease in sZmax at 400 “C in Fig. 1. Figure 3 shows the dependence of the microstructure factor R on the deposition temperature for samples grown at 30 Pa. The microstructure factor reaches a minimum at 350 “C at the same time as Q,,,._ exhibits a maximum at this temperature (Fig. 1). However, the hydrogen content for these samples (Fig. 4), determined by FTIR spectroscopy, decreases linearly from 11% to 5% when the deposition temperature rises from 200 “C to 400 “C without displaying any particular change at 350 “C. In Fig. 5, EB also decreases linearly from 2.02 eV to 1.88 eV as T, increases from 200 “C to 400 “C. The correlation between the hydrogen content and the optical gap energy E, is determined by comparing both Fig. 4 and Fig. 5. It is
0.3
143
extensively admitted that the optical gap energy is affected by hydrogen content [28], and the experimental data from the literature reveal a clear correlation between the optical gap energy and the bonded hydrogen density [29]. Von Roedern et al. [30] concluded, from photoemission experiments, that hydrogen incorporation removes states on the top of the valence band. However, Cody et al. [31] consider that the optical gap energy is determined by the degree of disorder in the material and that hydrogen indirectly affects the gap energy, releasing stress to the amorphous network. In addition to the saturation of dangling bonds, another important effect of hydrogen incorporation arises from its influence on the network cross-linking by lowering its self-coordination (Si-Si bonding). In effect, hydrogen introduces a fourth hetero coordination @i-H), which acts as a bonding terminal, thus leading to a looser network. The variations in the slope B and the energy &, of the Urbach edge with the deposition temperature are shown in Fig. 6 and Fig. 7 respectively. Both dependences
w 0.2 0.1 i
O.O ‘0 150
200 250 300 350 Deposition Temperature
400 (“C)
Fig. 3. Deposition temperature dependence of the microstructure parameter R, defined from the a-Si:H stretching modes. Deposition conditions are 30 Pa of pressure. 30 standard cm3 min-’ of SiH, flow and 8 W of r.f. power.
9
I
150
,
200 250 300 350 Deposition Temperature
400 (“C)
450
Fig. 6. Deposition temperature dependence of the slope B referred to the linear dependence of Ee,‘!2 on energy E of the Taut model. The deposition conditions were 30 Pa of pressure, 30 standard cm3 min- ’ of SiH, flow and 8 W of r.f. power.
k I”
41d 150
200 250 300 350 Deposition Temperature
400 (“C)
450
Fig. 4. Variation with the deposition temperature in the bonded hydrogen content in a-Si:H films, determined from the wagging mode of FTIR spectra. The deposition conditions are 30 Pa of pressure, 30 standard cm3 min _ ’ of SiH, flow and 8 W of r.f. power.
55
1
150
I
I
I
200 250 300 350 Deposition Temperature
I
400 (“C)
450
Fig. 7. Deposition temperature dependence of the Urbach edge energy, determined by PDS of a-Si:H thin films grown on quartz substrates at 30 Pa of pressure, 30 standard cm3 min-’ of SiH, flow and 8 W of r.f. power.
1.6 &I 150
200 250 300 350 Deposition Temperature
400 (“C)
450
Fig. 5. Optical gap energy dependence on deposition temperature of bulk a-Si:H. These values were calculated from dielectric function spectra, determined by elhpsometric measurements, taking into account the Taut model. The deposition conditions are 30 Pa of pressure, 30 standard cm3 min-’ of SiH, flow and 8 W of r.f. power.
indicate that the disorder in the amorphous network reaches a minimum around 300-350 “C and, furthermore, that disorder is higher at low deposition temperatures. The localized states density N, in the gap (Fig. 8), measured by PDS, also exhibits a minimum value in the interval 300-350 “C (N, = 3.7 x 1Or6 cmm3). However,
E. Bertrun et ul. / R.f: glow discharge u-Si
144
150
200 250 300 350 Deposition Temperature
400 (“C)
Fig. 8. Deposition temperature dependence of the density of localized states in the forbidden band of a-Si:H, determined by PDS measurements, The films of a-Si:H were deposited at 30 Pa of pressure. 30 standard cm3 min _ ’ of SiH, flow and 8 W of r.f. power.
the film deposited at 400 “C is the film that exhibits a higher density of states in the gap. The real-time ellipsometric trajectories corresponding to film growth have been analysed assuming a previously suggested model [ 181, where hemispherical nucleation, the coalescence phase and the homogeneous growth are taken into account. Figure 9 shows the film thickness at
:kGT r
Deposition
--330
300
temperature
-4;10 4504 (“C)
Fig, 9. Deposttion temperature dependence of the a-Si:H film thickness at the end of the nucleation and coalescence phases, determined from real-time ellipsometric measurements. The films were deposited at 30 Pa of pressure. 30 standard cm3 min 1of SiH, flow and 8 W of r.f. power.
the end of the coalescence phase as a function of deposition temperature. This parameter decreases to a minimum at a deposition temperature of 350 “C. This behaviour can be a consequence of two phenomena related to hydrogen content and surface hydrogen coverage. The nucleation phase depends on the state of the surface and of the hydrogen coverage, and therefore a delay in the final stage of the nucleation phase as a consequence of the decrease in hydrogen coverage above 350 “C can be expected. On the contrary, the coalescence phase can be completed earlier because of the increase in the dangling bonds between growing cylinders [ 181as the deposition temperature rises. Results in this paper reveal the influence of deposition temperature on different properties of a-Si:H. However, it is possible that this technological parameter does not directly affect the a-Si:H properties, but rather it acts on hydrogen through the deposition process mechanisms. At a deposition temperature of 350 “C, all studied properties exhibit an extreme behaviour. In spite of that, the characterization techniques used are independent of
each other, and the measured properties are based on different principles. This suggests that hydrogen incorporated in the network controls the a-Si:H properties. The important role of hydrogen in the structure of a-Si:H is well known. Hydrogen incorporation during the growth of a-Si:H thin films is controlled by two antagonistic kinetics. The first kinetics is associated with the flux of the impinging hydride radicals. The kinetics of the hydrogen contribution depends on the flux of the impinging radicals and on the number of hydrogen atoms associated with each type. The second kinetics which derives from two heterogeneous reactions is associated with the ejection of hydrogen from hydrogen-rich radicals (hydrogen desorption). These two kinetics compete severely, the resultant leading to a relatively small incorporation of hydrogen (about IO’:;) compared with the H:Si ratio of the impinging species (H:Si 2 300’;$ thus about 290”,, of the hydrogen has been ejected. Since the temperature is a determinant of the heterogeneous reaction. it allows us to “modulate” the hydrogen incorporation. Hydrogen incorporates in amorphous silicon under different bond configurations such as Si-H, Si-HZ. Si-H, and (Si-H,),. depending on preparation conditions. In the light of the results, everything seems to indicate a thermal activation of the surface mobility of adsorbed radical species on the growing film, limited by the desorption of hydrogen above 350 “C. The arguments supporting this assumption are stated as follows. (a) The material density (as a synonym of F~,,,_ in Fig. 1) increases as the deposition temperature increases before reaching the value of 350 ‘C. which suggests that the increase in surface mobility favours the formation of Si-H bonds with respect to other bonding‘forms. (b) The evolution of hydrogen content (Fig. 4) and the decrease in microstructure factor R which means a decrease in Si-H, and 2(Si-H) bonds with respect to the formation of Si-H bond (Fig. 3) are also indicative that films become denser. (c) The Urbach edge energy evolution in Fig. 7, which indicates a reduction in localized states in the band tails as the temperature rises to 350 ‘C, can be interpreted as a higher efficiency of the hydrogen in saturating dangling bonds. This interpretation is supported by the decrease in density of localized states as the temperature rises (Fig. 8). (d) The real-time elliljsometric results also indicate an increase in surface mobility, since both the nucleation phase and the coalescence phase become faster as the deposition temperature rises to 350 C.
4. Conclusion It becomes clear that the deposition temperature is a critical parameter which has beneficial effects up to
E. Bertran et al. / R.f. glow discharge a-Si
350 “C by lowering the microstructure parameter, the material density, the density of localized states and the Urbach edge as the deposition temperature increases. These dependences indicate an increase in surface mobility which favours the formation of Si-H bonds with respect to other bonding forms and a higher efficiency of the hydrogen in saturating dangling bonds. Above this threshold temperature, at which the desorption of hydrogen incorporated in the network becomes important, the density of states, the Urbach edge, the material density and the microstructure parameter increase. The main mechanism active in the growth process is the thermal activation of the surface mobility of adsorbed radical species on the growing surface, which are limited by the hydrogen surface desorption. Above a deposition temperature of 350 “C, less hydrogen coverage occurs during the a-Si:H deposition process.
6 7 8 9 10 11 12 13 14 15
16 17
Acknowledgments
18
The authors acknowledge Professor J. L. Morenza and Dr. J. Andreu for their suggestions and they thank Dr. A. Lloret for fruitful discussions. This work has been supported by the CAICYT de1 Ministerio de Education y Ciencia of Spain under Contract 798/84 and the CICYT of Spain under Contract MAT 955190.
19 20
21 22 23
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