Structural and optical properties of hydrogenated amorphous silicon ...

PHILOSOPHICAL MAGAZINE B, 2002, VOL. 82, NO. 1, 35±46

Structural and optical properties of hydrogenated amorphous silicon±carbon alloys grown by plasmaenhanced chemical vapour deposition at various rf powers G. Ambrosone, U. Coscia Dipartimento di Scienze Fisiche ed UnitaÁ Istituto Nazionale per la Fisica della Materia, UniversitaÁ di Napoli `Federico II’, Via Cintia, 80125 Napoli, Italy

S. Ferrero, F. Giorgis, P. Mandracci and C. F. Pirriy Dipartimento di Fisica ed UnitaÁ Istituto Nazionale per la Fisica della Materia, Politecnico di Torino, Corso Duca degli Abruzzi 24, 10129 Torino, Italy [Received 18 May 2001 and accepted 23 May 2001]

Abstract Hydrogenated amorphous silicon±carbon alloys (a-Si1 x Cx : H† deposited by plasma techniques are important for a large variety of applications from microelectronics and optoelectronics to microelectromechanic and photonic structures. However, it is known that as the carbon concentration increases in the alloy, the properties of the ®lms strongly depend on the deposition technique and growth conditions. The aim of the present work is the investigation of structural and optical properties and defects of a-Si1 x Cx : H ®lms grown by plasma-enhanced chemical vapour deposition in a silane±methane plasma at rf power densities up to 60 mW cm 2 and with carbon contents x up to 0.6. Films have been produced and characterized with low defect densities, below 1:3 1017 cm 3 , deposition rates higher than 0.1 nm s 1 , evidence of chemical order for a near-stoichiometric region and high radiative e ciency. The results have been discussed in the light of consolidated models of the deposition chemistry, structure and physical properties of a-Si1 x Cx : H.

} 1. Introduction Hydrogenated amorphous silicon±carbon alloys (a-Si1 x Cx : H) deposited by plasma techniques have been used in a large variety of microelectronic and optoelectronic applications (Kanicki 1991). This is because, by changing the material composition, it is possible to obtain ®lms with optical gaps ranging from 1.9 up to 3.6 eV and with high-e ciency room-temperatur e luminescence. Recently, other promising applications of a-Si1 x Cx : H have been proposed in microelectromechanic systems (Sarro 2000) and in photonic structures (Giorgis 2000). However, aSi1 x Cx : H materials have physical properties strongly dependent on the deposition technique and growth conditions, in particular for carbon contents approaching or exceeding stoichiometry. The aim of the present work is the investigation of structural, optical and radiative recombination properties of a-Si1 x Cx : H ®lms grown by y Email: [email protected]

Philosophica l Magazine B ISSN 1364±2812 print/ISSN 1463±6417 online # 2002 Taylor & Francis Ltd http://www.tandf.co.uk/journals DOI: 10.1080 /1364281011006927 9

36

G. Ambrosone et al.

plasma-enhanced chemical vapour deposition (PECVD) from a silane (SiH4 † methane (CH4 ) plasma at various rf powers. Films with carbon contents x ranging from 0.1 to 0.6 have been produced and investigated with respect to, ®rstly, their optical and defect properties by transmittance±re¯ectance spectroscopy and photothermal de¯ection spectroscopy (PDS), secondly, their structural and compositional properties by infrared (IR) spectroscopy, Rutherford back-scattering spectroscopy (RBS) and elastic recoil detection analysis (ERDA) and, thirdly, radiative properties by room-temperatur e continuous-wave (CW) photoluminescence (PL) spectroscopy. Low defect densities, high radiative e ciencies and evidence of chemical order near stoichiometry have been obtained for rf powers up to 60 m W cm 2 . } 2. Experimental details a-Si 1 x Cx : H samples were deposited by standard PECVD at 13.56 MHz on Si (100) wafers, Corning 7059 glass and quartz substrates. For all the samples the deposition conditions have been chosen so as to optimize their properties as reported elsewhere (Rava et al. 1996, Giorgis et al. 1997, 1998b). The CH4 percentage in the SiH 4 ±CH4 plasma was varied between 50 and 98% with a total ¯ow rate of 40 sccm in order to keep the dwell time of the molecule constant at 0.8 s. The rf power was varied between 10 and 32 W, corresponding to power densities in the range from 10 to 60 mW cm 2 . The other deposition conditions were ®xed as follows: substrate temperature, 200°C; pressure, 0.35 Torr; electrode distance, 15 mm. The elemental silicon and carbon contents of the ®lms were obtained by RBS (Chu et al. 1978). The absorption coe cient above 10 4 cm 1 was obtained by transmittance±re¯ectance spectroscopy with a Perkin±Elmer Lambda 9 spectrophotometer, operating in the range 200±3200 nm, and the absorption coe cient below 104 cm 1 from PDS by a conventional experimental set-up (Jackson et al. 1981). The optical gap was evaluated as E04 , that is the energy at which the absorption coe cient is 104 cm 1 . The gap defect densities and the states at the valence-band edge were evaluated through the excess in absorption coe cient for energies below the Urbach region, for all the samples. The ®lm structure was investigated by IR spectroscopy with a Perkin±Elmer 2000 Fourier transform IR apparatus. The spectra were recorded in absorption mode in the range 400±4000 cm 1 with a resolution of 1 cm 1 . CW PL measurements were performed at room temperature using an argon laser or a Xe±Hg lamp as exciting sources (lines at 2.71 and 3.4 eV respectively), and a crystalline silicon photodiode whose signal was processed by a digital lock-in ampli®er. } 3. Results and discussion Figure 1 shows the carbon content x in a-Si1 x Cx : H ®lms as a function of CH4 percentage in the plasma for the sets deposited at various rf powers. We observe that, on varying the methane percentage (CH4 %) from 50% to 98%, the carbon incorporation ranges from 0.13 to 0.55 for samples deposited at a low rf power (10 m W cm 2 ) and from 0.20 to 0.60 for samples deposited at a high rf power (60 m W cm 2 ). The trend of carbon incorporation seems typical of the so-called low-power regime (Schmidt et al. 1985, Bullot et al. 1987), as con®rmed by the decrease in the ®lm deposition rate as the carbon content increases (reported in the inset of ®gure 1). However, we observe a few novel results.

Structural and optical properties of a-Si1 x Cx : H

37

(i) The carbon incorporation for our samples exceeds 0.5, which is considered the upper limit for low-power regime conditions. (ii) The carbon incorporation and deposition rate are in¯uenced by changing the rf power. These results suggest, as already con®rmed by Bullot and Schmidt (1987), Rava et al. (1996) and Pereyra et al. (1998), that SiH 4 ±CH4 plasmas are still far from completely understood and that the `low-power’ and the `high-power’ regimes do not exhaustively describe all the a-Si1 x Cx : H growth mechanisms. This is very important for the growth of wide-bandgap a-Si1 x Cx : H ®lms with optoelectronic properties suitable for device applications. In fact, device-quality a-Si1 x Cx : H has been produced only in the `low-power regime’ with x < 0:5 and deposition rate below 0.01 nm s 1 . The `high-power regime’, which grows samples at carbon contents exceeding 0.5, thereby covering a wider range of optical properties, and deposition rates higher than 0.2 nm s 1 , gives rise to materials with poor optoelectronic properties as a consequence of plasma conditions. For the a-Si1 x Cx : H ®lms of the present paper the deposition conditions are not far from those of the `lowpower regime’ (in fact the rf power density ranges from 10 to 60 mW cm 2 and the total ¯ow rate is 40 sccm); the deposition rates are in¯uenced by the rf power and can be increased to 0.05 nm s 1 for samples having x > 0:5. Figures 2 and 3 show the optical and defect properties (as deduced from absorption coe cient) respectively of the a-Si1 x Cx : H ®lms deposited at two di erent rf powers as functions of the carbon content. The optical gap E04 (®gure 2 (a)) covers a range from 2.0 up to 3.3 eV, (®gure 2 (b)), increasing linearly as a function of the carbon content x in ®lms and

Figure 1. Carbon content x in a-Si1 x Cx : H ®lms grown by PECVD at two di erent rf power densities of 10 mW cm 2 (low power) and 60 mW cm 2 (high power) in SiH4 ±CH 4 plasma as a function of CH4 fraction (the inset reports the deposition rate as a function of the carbon content x in the ®lms).

38

G. Ambrosone et al.

Figure 2. (a) Optical gap E04 and (b) refractive index at 1500 nm for a-Si1 x Cx : H ®lms grown by PECVD at two di erent rf power densities of 10 mW cm 2 (low power) and 60 mW cm 2 (high power) in SiH4 ±CH4 plasma as a function of the carbon content x.

the refractive index in the range 1.9±3.3 decreases linearly as a function of the carbon fraction. The Urbach energy EU (®gure 3 (a)) increases as the carbon content in the ®lms increases with di erent trends for x < 0:5 and for x > 0:5 and without signi®cant e ects of rf power. For x < 0:5 the Urbach energy, starting from 75 meV reaches values around 150 meV; for x > 0:5, EU increases to 300 meV. The defect density (®gure 3 (b)) increases starting from values of around 2 1016 cm 3 for x ˆ 0:13 up to values close to 1:3 1017 cm 3 for x > 0:5. It is interesting to note, by comparing the present results with the literature (Giorgis et al. 1998c), that the reported defect density is very low even for samples having x > 0:5 and an optical gap of around 3.3 eV. This is very important for applications of wide-bandgap aSi1 x Cx : H ®lm in optoelectronic devices. The PL spectra have been collected at room temperature with excitation over the optical gap E04 in order to avoid the typical variations in emission when the excitation energy is below E04 (Tessler and Solomon 1995). The experimental data exhibit a Gaussian shape, and three systematic trends can be observed. (i) The PL spectra shift towards higher energies. (ii) The spectra widths increase. (iii) The e ciency at room temperature increases as the carbon content increases.


39

Figure 3. (a) Urbach energy and (b) defect density ND for a-Si1 x Cx : H ®lms grown by PECVD at two di erent rf power densities of 10 mW cm 2 (low power) (&) and 60 mW cm 2 (high power) (&) in SiH4 ±CH4 plasma as a function of the carbon content x.

Figures 4 (a) and (b) report the PL e ciency and the energy EPL of maximum emission respectively as functions of the optical E04 gap for both sets of ®lms. The PL e ciency versus E04 gap (®gure 4 (a)) shows an increase of more than three orders of magnitude as E04 varies from 2.0 to 2.8 eV, corresponding to a carbon content from 0.13 to 0.50; afterwards a saturation at E04 ˆ 2:8 eV is reached for x > 0:5. The energy position EPL of the maximum of the emission spectra monotonically shifts from 1.5 to 1.9 eV as E04 increases from 2.2 to 3.3 eV (®gure 4 (b)). No di erences have been evidenced in the emission spectra for the sets deposited at di erent RF power densities, supporting the idea that, in our growth regime (even if it is not fully observing the `low-power regime’ rule), physical properties are controlled only by carbon incorporation. The inset of ®gure 4 (a) shows the Stokes shift, which is the di erence between the E04 gap and the energy position EPL of the maximum of the emission spectra, as a function of Urbach energy. It is possible to observe two regions. (i) For an Urbach energy below 200 meV, that is for x < 0:5, the emission data can be interpreted in terms of the static disorder model (Dunstan and

G. Ambrosone et al.

40

Figure 4. (a) Photoluminescence e ciency and (b) energy position of the maximum emission EPL for a-Si1 x Cx : H ®lms grown by PECVD at two di erent rf power densities of 10 mW cm 2 (low power) (&) and 60 mW cm 2 (high power) (&) in SiH4 ±CH 4 plasma as a function of the optical gap E04 (the inset of (a) shows the di erence between E04 and EPL as a function of the Urbach energy for the same ®lms).

Boulitrop 1984, Giorgis et al. 2000), which attributes the PL spectra to recombination at band-tail states. Following this model, the increase in PL ef®ciency as a function of carbon content can be due to the decrease in charge mobility, and the Stokes shift E04 EPL is given by E04

EPL ˆ EU ln …43 pR3C N0 EU †

…1†

where RC is the carrier distance at which radiative and non-radiative rates are equal and N0 is the density of states at the band edge (Giorgis et al. 2000). If we assume N 0 ˆ 5 10 21 cm 3 eV 1 , the ®t of the data plotted in the inset of ®gure 4 (a) for Urbach energies below 200 meV (shown in ®gure 4 (a) as a solid line) gives RC ˆ 10 nm, which is in agreement with results obtained by Boulitrop and Dunstan (1983) for amorphous silicon. (ii) For Urbach energies above 200 meV, that is x > 0:5, the carbon clusters favour the formation of sp2 bonds. The presence of sp2 clusters adds another radiative tail to tail recombination path (Giorgis et al. 1998a). For this


41

reason, in carbon-rich samples, we observe a saturation both in PL ef®ciency versus E04 and in the E04 EPL shift versus EU : In order to understand the optical, defect and luminescence properties, IR spectroscopy has been applied to all the samples. The IR spectra of the a-Si1 x Cx : H deposited ®lms, can be divided into ®ve regions (Demichelis and Pirri 1995). (I)

(II)

(III)

(IV) (V)

The region between 2800 and 3100 cm 1 is due to CHn stretching modes in both sp3 and sp2 con®gurations. In ®lms grown from SiH4 ±CH4 gas mixtures there are two resolved peaks; in the range 2920±2960 cm 1 the vibrations can be attributed to asymmetric stretching of CH3 groups in the sp3 con®guration (2960 cm 1 ) and the asymmetric stretching of sp3 CH2 groups (2920 cm 1 ); in the range 2860±2880 cm 1 the modes can be attributed to sp3 CH2 symmetric vibrations (2860 cm 1 ) and sp3 CH3 symmetric and/or CH2 stretching vibrations (2880 cm 1 †. The mode at 2100 cm 1 is attributed to the stretching vibrations of single SiÐH bonds with one, two or three carbon atoms attached to a silicon atom or to a shift of the SiH stretching mode at 2000 cm 1 towards 2100 cm 1 due to the formation of SiHn groups. The region between 1200 and 1500 cm 1 is attributed to the CHn bending or scissoring modes. The 1250 cm 1 peak is due to the symmetric bending mode of CH3 attached to the silicon atom. The doublet at 1350 and 1404 cm 1 can be attributed to SiÐCH3 asymmetric bending vibrations, or CH2 wagging (1350 cm 1 † and CH2 scissoring or bending vibrations (1400 cm 1 ). The region between 950 and 1100 cm 1 is assigned to CHn rocking or wagging vibrations. The mode at 670 cm 1 is attributed to SiH wagging vibration and the strong feature at 780 cm 1 is assigned to the SiÐCH3 rocking or wagging mode or to the SiÐC stretching mode.

No signal has been detected in the region between 1500 and 1600 cm 1 owing to CÐC stretching vibrations, which is usually attributed to the presence of large sp2 carbon clusters. It is interesting to note that no di erences have been evidenced in the spectra of a-Si1 x Cx : H ®lms deposited at di erent rf power densities, except for that due to the di erent carbon amounts. The integrated intensity of the IR peaks is related to the bond concentrations. The integrated absorption is given by the relationship Iˆ

…

¬…!† d!; !

…2†

where ¬…!) is the IR absorption coe cient; the integration is performed over the deconvoluted IR peaks after subtraction of the baseline. Some attempts have been made to perform an absolute evaluation of bond concentration in a-Si1 x Cx : H ®lms by IR spectroscopy in order to obtain information on material structure. The concentration of oscillators is given by NXY ˆ AXY

…

¬…!† d!; !

XY ˆ SiC; SiH or CH;

…3†

42

G. Ambrosone et al.

where ¬…!† is the absorption coe cient of the 800 cm 1 band for SiÐC vibrations, the 2000±2100 cm 1 band for SiÐH vibrations and the 2800±3100 cm 1 band for CÐH vibrations. Several determinations of the ASi C , ASi H and AC H constants have been performed. It was calculated that ASi C ˆ 3 10 19 cm 2 (Bullot and Schmidt 1987, Basa and Smith 1990, Van Swaaij et al. 1994, Friessnegg et al. 1998, Pereyra et al. 1998), ASi H ˆ 1:4 10 20 cm 2 (Lanford and Rand 1978, Mui et al. 1987, Basa and Smith 1990, Friessnegg et al. 1998, Pereyra et al. 1998) and AC H ˆ 1:35 10 21 cm 2 (Guivarach et al. 1980, Friessnegg et al. 1998, Pereyra et al. 1998). Figures 5 (a) and (b) show SiH and CH bond concentrations respectively, deduced from equation (3), as a function of the carbon content in the ®lms. SiH bond concentrations exhibit values around 2 10 21 cm 3 up to a maximum of 6 10 21 cm 3 for x in the range from 0.13 to 0.50; afterwards the concentration decreases to 3:5 10 21 cm 3 for x ˆ 0:6 (®gure 5 (a)). The CH bond concentration monotonically increases as function of carbon content from 1 10 21 up to

Figure 5. (a) SiH and (b) CH bond concentrations, as deduced from IR spectra, for aSi1 x Cx : H ®lms grown by PECVD at two di erent rf power densities of 10 mW cm 2 (low power) (&) and 60 mW cm 2 (high power) (&) in SiH4 ±CH 4 plasma as a function of the carbon content x.


43

1:5 1022 cm 3 . IR spectra show that SiH n groups are mainly in the monohydride form (n ˆ 1) and only a small fraction in the dihydride con®guration (n ˆ 2), as evidenced by the presence of the 845±890 cm 1 doublet. Moreover CHn groups are prevalently in the dihydride (n ˆ 2) and trihydride (n ˆ 3) form, favouring the microvoid formations in the tetrahedral SiÐSi and SiÐC network with the consequent creation of defects and disorder. In fact, ®gures 6 (a) and (b) report the defect density ND and the Urbach energy (usually considered a measure of disorder in the covalent network) respectively as functions of the XÐH ˆ ‰SiÐHŠ ‡ ‰CÐHŠ group concentrations. Figure 6 (a) shows an almost linear correlation between defect densities and XÐH group concentrations. Figure 6 (b) suggests that, for carbon concentrations below 0.5 (under-stoichiometri c samples), also disorder is linearly correlated to the concentration of XÐH groups. At carbon contents close to stoichiometry (around x ˆ 0:5), the Urbach energy becomes independent of the XÐH concentration, which can be explained by the presence of other contributions to disorder related to microstructure.

Figure 6. (a) Defect density ND and (b) Urbach energy for a-Si1 x Cx : H ®lms grown by PECVD at two di erent rf power densities of 10 mW cm 2 (low power) (&) and 60 mW cm 2 (high power) (&) in SiH4 ±CH4 plasma as a function of XÐH ˆ ‰SiÐHŠ ‡ ‰CÐHŠ bond concentration deduced by IR spectra.

44

G. Ambrosone et al.

Figure 7 (a) shows the total hydrogen content, obtained by IR spectroscopy, as a function of carbon content for the a-Si1 x Cx : H ®lms deposited at two di erent rf power densities. We can observe a monotonic increase in the hydrogen concentration from 0:3 10 22 cm 3 for x ˆ 0:13 up to 2 1022 cm 3 for x ˆ 0:6. By using the silicon and carbon atom doses, as obtained by RBS, the hydrogen atomic percentage has been computed. It is possible to observe, as reported in ®gure 4 (a), that the total hydrogen ranges from 12 at.% in silicon-rich ®lms up to 32 at.% in carbon-rich ®lms. Joining together the results of ®gures 5 and 7 (a) we can conclude that, for samples having a carbon content x < 0:4, the hydrogen is incorporated at almost the same amount in SiÐH and CÐH groups while, for those with higher carbon contents, hydrogen is mainly incorporated in CÐH groups. Figure 7 (b) reports the integrated intensities (following equation (2)) of IR vibrations at 780±800 cm 1 as a function of carbon content (for a-Si1 x Cx : H ®lms

Figure 7. (a) Hydrogen concentration as deduced from IR spectra, and (b) integrated intensities of IR peaks at 780 cm 1 for a-Si1 x Cx : H ®lms grown by PECVD at two di erent rf power densities of 10 mW cm 2 (low power) and 60 mW cm 2 (high power) in SiH4 ±CH4 plasma as a function of the carbon content x.


45

deposited at two di erent rf power densities). In the same ®gure the solid curve is the theoretical trend of SiC bond concentration computed through RBS and IR results (Robertson 1992, Rovira and Alvarez 1997), by supposing complete chemical order in the network. In these conditions, taking into account the SiÐH and CÐH bonds, the complete presence of SiÐC bonds in the network structure is reached for x ˆ 0:55. Two important features can be noted in ®gure 7 (b): ®rstly, for x < 0:4 the concentration of SiÐC bonds is lower than that corresponding to complete chemical order; secondly for x > 0:4 the experimental data ®t the trend of complete chemical order up to the maximum carbon content in deposited samples (x ˆ 0:6). Finally the maximum of SiÐC bond concentration is reached at x ˆ 0:55 corresponding to a network completely formed by SiÐC, CÐH and SiÐH groups. } 4. Conclusions a-Si 1 x Cx : H samples deposited at rf power densities in the range 10± 60 mW cm 2 by PECVD in SiH4 ±CH4 gas mixtures have been investigated with respect to their optical, structural, defective and radiative recombination properties for carbon contents up to 0.6. The main results obtained can be summarized as follows. (1) a-Si1 x Cx : H ®lms with carbon contents x up to 0.6 can be obtained in the low-power regime (rf power densities below 100 mW cm 2 † at deposition rates above 0.1 nm s 1 . (2) The ®lms show an optical gap in the range 1.9±3.3 eV and exhibit a high ef®ciency of room-temperatur e radiative recombination. (3) The structure of the ®lms investigated by IR spectroscopy, RBS and ERDA has evidenced the possibility of growing chemically ordered amorphous a-Si1 x Cx : H ®lms when the composition is close to stoichiometry. (4) Finally, the defect density in the gap, evaluated by the PDS technique, is below 1:3 10 17 cm 3 for all the compositional range. References Basa, D. K., and Smith, F. W., 1990, Thin Solid Films, 192, 121. Boulitrop, F., and Dubnstan, D. J., 1983, Phys. Rev. B, 28, 5923. Bullot, J., and Schmidt, M. P., 1987, Phys. Stat. sol. (b), 143, 345. Chu, W. K., Mayer, J. M., and Nicolet, M. A., 1978, Backscattering Spectrometry (New York: Academic Press.). Demichelis, F., and Pirri, C. F., 1995, Solid State Phenom., 44±46, 385. Dunstan, D. S., and Boulitrop, F., 1984, Phys. Rev. B, 30, 5945. Friessnegg, T., Boudreau, M., Mascher, P., Knights, M., Simpson, P. J., and Puff, W., 1998, J. appl. Phys., 84, 786. Giorgis, F., 2000, Appl. Phys. Lett., 77, 522. Giorgis, F., Giuliani, F., Pirri, C. F., Tagliaferro, A., and Tresso, E., 1998a, Appl. Phys. Lett., 72, 2520. Giorgis, F., Giuliani, F., Pirri, C. F., Tresso, Conde, J. P., and Chu, V., 1998b, J. noncrystalline Solids, 227±230, 465. Giorgis, F., Giuliani, F., Pirri, C. F., Tresso, E., and Coscia, U., 1998c., Amorphous Silicon and its Alloys, emis Data Reviews Series 19, edited by T. Searle (Exeter: INSPEC IEE Publication, Short Run Press). Giorgis, F., Mandracci, P., Dal Negro, L., Mazzoleni, C., and Pavesi, L., 2000, J. noncrystalline Solids, 266±269, 588. Giorgis, F., Pirri, C. F., Tresso, E., Rigato, V., Zandolin, S., and Rava, P., 1997, Physica B, 229, 490.

46


Guivrach, A., Richard, J., Le Contellec, M., Ligeaon, E., and Fontenille, J., 1980, J. appl. Phys., 51, 2167. Jackson, W. B., Amer, N. M., Boccara, C., and Fournier, D., 1981, Appl. Optics, 20, 1333. Kanicki, J., 1991, Amorphous and Microcrystalline Semiconductor devices (London: Artech House). Landford, W. A., and Rand, M. J., 1978, J. appl. Phys., 48, 2474. Mui, K., Basa, D. K., and Smith, F. W., 1987, Phys. Rev. B, 25, 8089. Pereyra, I., Carreno, M. N. P., Tabacniks, M. H., Prado, R. J., and Fantini, M. C. A., 1998, J. appl. Phys., 84, 2371. Rava, P., Crivini, G., Demichelis, F., Gorgis, F., and Pirri, C. F., 1996, J. appl. Phys., 80, 4116. Robertson, J., 1992, Phil. Mag. B, 66, 615. Rovira, P. I., and Alvarez, F., 1997, Phys. Rev. B, 55, 4426. Sarro, P. M., 2000, Sensors Actuators, 82, 210. Schmidt, M. P., Solomon, I., Tran-Quoc, H., and Bullot, J., 1985, J. non-crystalline Solids, 77±78, 849. Tessler, L. R., and Solomon, I., 1995, Phys. Rev. B, 52, 10 952. Van Swaaij, R., Bernstein, A. J. M., Van Sark, W. G., Herremans, H., Bezemer, J., and Van der Weg, W. F., 1994, J. appl. Phys., 76, 251.