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Boron and phosphorus doping of a-SiC:H thin films by means of ion implantation
ELSEVIER
Thin Solid Films 265
( 1995) 113-l 18
Boron and phosphorus doping of a-SiC:H thin films by means of ion implantation F. Demichelis a, G. Crovini a, C.F. Pirri a, E. Tress0 a, R. Galloni b, C. Summonte b, R. Rizzoli b, F. Zignani ‘, P. Rava d aDipartimento di Fisica, Politecnico di Torino, Italy b CNR-Lamel, Bologna, Italy ’ Facoltci di Ingegneria. Vniversitri di Bologna, Italy d Elettrorava S.p.a.. Savonera, Torino, Italy Received 21 November
1994; accepted 28 March 1995
Abstract A detailed study of the structural and optoelectronic properties of boron- and phosphorus-implanted a-SiC:H thin films is presented. The films have been deposited by ultra high vacuum plasma-enhanced chemical vapour deposition and have energy gap values of about 2 eV. The effects of varying carbon content, boron-and phosphorus-implanted dose and annealing temperatures are reported. It is found that ion implantation strongly modifies the structure, producing an increase of monohydride bonds. After annealing the defect density is mainly affected by dopant type and concentration rather than by residual radiation damage. We also show that no enhanced gap shrinkage is produced by boron with respect to phosphorus doping if the film composition is the same. A comparison with gas phase doping is also reported. Keywork
Amorphous
materials;
Ion implantation;
Silicon carbide; Phosphorus;
1. Introduction Intrinsic a-SiC:H films are widely used for solar cells, lightemitting diodes, photodiodes, sensors and several other applications [ l-31, taking advantage of the possibility of tailoring their optical gap in the range 1.8-3.0 eV. Thin layers of aSiC:H with high conductivity, high transparency and low defect density are needed for instance in solar cell technology and can be obtained by gas-phase doping with either boron (p-type material) or phosphorus (n-type material) [4,5]. Generally the p-type doping is produced by mixing into the SiH4 gaseous phase B2Hs, while PH3 is used for n-type doping. The properties of a-SiC:H films doped using gases such as B,H, and PH, have been widely investigated [ 6,7]. Post-deposition doping techniques such as ion implantation [ 8-121 offer the advantage of a better control of the impurity concentration and the possibility of creating doped patterns in an undoped matrix. Nevertheless the creation of a great deal of damage due to the implantation process is the major drawback of the ion-implantation technique. Both point defects and voids are originated by the ion bombardment, resulting in a significant increase of the density of localized gap states. A feasible method to recover the damage is to perform an annealing process [ 13-151. 0040~6090/95/$09.50 0 1995 Elsevier Science S.A. All rights reserved SSDIOO40-6090(95)06618-7
Boron
Data about high-dose ion implantation doping on amorphous silicon carbide alloy films with energy gap values in the range 1.94-2.10 eV have already been published [ 16,171. In this paper we present a detailed study on the structural and optoelectronic properties of a-SiC:H thin films deposited by ultra high vacuum plasma-enhanced chemical vapour deposition (UHV-PECVD), implanted either with boron or phosphorus. The evolution of bonding structure, electrical activation and residual damage as a function of carbon content, implanted concentration and annealing temperature are reported.
2. Experimental Samples were deposited by UHV-PECVD in a SiH4 + CH, gas mixture. The deposition conditions are reported in Table 1: they were optimized in order to have device-quality a-SiC:H films with an energy gap of about 2 eV. In order to obtain homogeneous doping through the sample thickness (about 0.3 p.m for all the samples), multiple energy phosphorus- and boron-ion implantation was performed. Doping uniformity better than 5% was obtained by successively implanting at three energies, which were 15, 55, 140
114
F. Demichelis et al. /Thin Solid Films 265 (1995) 113-118
Table 1 Deposition conditions Substrate temperature Total pressure R.f. power CH, flow rate SiH, flow rate
175-190 “C 80 Pa 2.8 lo-’ W cm-* 40-70 seem 40 seem
keV for P ions and 15, 33, 60 keV for B ions. The total implanted peak concentration ranged from lOI to 5 X 10” cmm3. In order to study the recovery of radiation damage after ion implantation, samples were isochronally annealed for 1 h in air at increasing temperatures between 150 and 400 “C. The optical absorption coefficient was obtained by transmittance and reflectance spectroscopy in the fundamental absorption region by means of a Perkin-Elmer UV-Vis-NIR Lambda 9 spectrophotometer in the 0.2-1.5 pm range and by photothermal deflection spectroscopy (PDS) in the subgap region. The nature of chemical bonding was studied by infrared spectroscopy in the absorption mode using a Perkin-Elmer FTIR 2000 spectrophotometer between 400 and 4000 cm- ’ The absorption coefficient was deduced from the transmission spectra taking into account the influence of interference fringes and absorption due to the silicon substrate. The deconvoluted peaks corresponding to the different vibrational modes and their integrated intensities were determined. Steady-state dark conductivity at 27 “C and conductivity versus temperature were measured in a planar configuration in a MMR vacuum chamber equipped with a K20A temperature controller. Different substrates suitable for different characterizations were used: Corning glass 7059 for electrical measurements, optically polished fused silica for optical measurements and PDS and FZ crystalline silicon for infrared (IR) spectroscopy. 3. Resultsand discussion 3.1. Structural properties The IR absorption spectra have been measured on as deposited, as implanted and annealed films with different CH, flow rates. All the spectra show the characteristic vibrational modes typical of a-SiC:H with a low carbon content [ 181 No significant variation as a function of carbon concentration has been observed in the spectra of as-deposited films. In Fig. 1 the deconvoluted absorption spectra, after subtraction of the interference fringes, are reported for the film deposited with 40 seem of Cl& flow rate, and, for the same film, after implantation with boron at different concentrations and after annealing at temperatures between 250 and 350 “C. The arrows indicate the modes of vibration, which are reported in
600
1006'
600
0
2000
2200
[cm-‘]
Fig. 1. Infrared absorption spectra of a-SiC:H films implanted with boron at two different concentrations and annealed at increasing temperature.
Table 2. A similar trend was obtained for the same film when it was implanted with phosphorus. By comparing the spectra of implanted films with those obtained on as-deposited ones it is evident that: there is a broadening of the peak at 640 cm-‘, due to SiH rocking and a decrease of the 780 cm-’ due to the SiCH3 rocking/wagging mode; - there is a disappearance of the doublet structure at 845890 cm-‘, attributed to (SiH,) n scissors modes; there is a decrease of the stretching modes of SiH2 (2070 cm-‘) and a widening of the 2000 cm-’ peak, due to SiH stretching modes. Doping by ion implantation is associated with induced structural changes: in fact the observed trends in the IR spectra depend on the implanted concentrations and are more marked in the heavily implanted films. Ion implantation promotes monohydride bonds at the expense of the initially dominating SiH2 and Six-H vibrations and modifies the characteristic microstructure of the a-SiC:H network. We note that these variations in IR spectra are attributed to radiation damage, as dopant concentrations similar to those used in this paper, when introduced by gas phase doping, do not influence the a-SiC:H IR spectrum [ 191.
Table 2 IR vibrational peak assignments 1 2 3 4 5 6
640 cm-’ 780 cm-’ 845-890cm-’
1000 cm-’ 2010 cm-’ 2070 cm-’
SiH wagging and/or rocking Si-CH3 and/or Sic rocking/wagging SiH, scissors CH, wagging SiH stretching SiH2 stretching
F. Demichelis etal. /Thin Solid Films 265 (1995) 113-118
B imp. 5 10zD cmS3
7 3-
I
Pimp. 1P
\ \
‘2070 .
.\\
.i ‘2wo
._/\/
cm4
3
‘-‘\
2 -’
P imp. 1018 cm-3
daesp. imp. as
250
300
350 400 T
‘annealing
c&
p,“p.
Boron
250
300 350
400
i3
I
[ “‘1
Fig. 2. Integrated absorption intensity of SiH and SiHz stretching peaks for films deposited with 40 seem CH., flow rate and implanted with boron and phosphorus at two different concentrations.
In order to study the radiation damage and its recovery, the integrated intensity, defined as: ff( a) do w
In=]---
has been evaluated for the absorption peaks centred at 20102020 cm- ’ and at 2070 cm- ’ at different annealing temperatures for the film deposited with 40 seem of methane flow and implanted at different boron and phosphorus concentrations. These are reported in Fig. 2. It is evident that the radiation damage is stronger for the higher implanted concentrations, for which the increase of the SiH integrated intensities is so strong to overcome the value of the SiH, one. When the samples are annealed at temperature between 250 and 300 “C a slight decrease of the SiH peak is accompanied by a slight increase of the SiH* one, indicating a tendency towards a structural recovery of the radiation damage. For annealing temperatures higher than 300 “C the decrease of the H-related peaks clearly indicates, in all the films, the onset of hydrogen evolution which, above 400 “C, induces stress and peeling-off of the samples.
13 t
I
I
I
as dep.
as imp.
150
I
200
I
250
I,,
,
300
as as dep. imp.
T annealing
150
.(
200
250
300
[“Cl
Fig. 3. Taut’s energy gaps versus annealing temperature of a-SiC:H films deposited with 40 seem CH, flow rate, implanted with boron and phosphorus at varying concentrations.
maximum value, whereas for higher temperatures the decrease of Eg is due to hydrogen evolution. In Fig. 4, the Eg values are reported for the samples deposited at CH, flow rates of 40 and 70 seem as a function of the annealing temperatures and for two different P and B concentrations. For dopant concentrations as low as lo’* cme3 the energy gap depends on carbon content. For higher implanted concentrations, after annealing at 250 “C, the Eg values lie in the range 1.8-l .9 eV, independent of the CI-L,flow rate. We would like to underline one feature that comes out from our results. As will be shown in the following (see Section 3.3), the doping is electrically effective for concentrations higher than about lOi cmd3. For those concentrations, and after annealing at 250 “C, the gap shrinkage in the case of B or P doping by ion implantation, is approximately the same. In the case of doping by the gas phase, instead, it was previously reported [20] that boron doping is accompanied by a more enhanced gap
3.2. Optical properties In Fig. 3 the optical gap Eg, evaluated from the Taut plot, is reported as a function of annealing temperature for the sample deposited at a CH4 flow rate of 40 seem implanted at four different B and P concentrations. The radiation damage produced by ion implantation results in a gap shrinkage which is more marked for higher implanted concentrations. The amount of damage also depends on the atomic number of the implanted ions: the gap shrinkage detectable, in the case of phosphorus implantation, even at low concentration, is a consequence of the larger weight of P with respect to B ion. The implantation damage is partially recovered by annealing. For each concentration, at annealing temperatures between 250 and 275 “C the Taut gap of implanted films reaches the
:1:1.:_::.1.‘: 1yL~z= I%. Ep. 150
200
250
300
&
Tannealing
#&
150
200
250
300
[ “Cl
Fig. 4. Taut’s energy gaps versus annealing temperature for a-SiC:H films deposited with different CH, flow rates and implanted at two different boron and phosphorus concentrations.
116
F. Demichelis et al. /Thin Solid Films 265 (1995) 113-l 18
the presence of boron atoms rather than to the implantation process. 3.3. Electrical properties
I. 5.
i&.
. 200
. 300
I
400
I. dt;.
,, la,6.
200
, 300
I 400
T annealing [“Cl Fig. 5. Urbach energy and defect densities of films deposited with 40 seem of CH, flow rate and implanted with boron and phosphorus at two different concentrations.
shrinkage than phosphorus doping. In Ref. [ 201 and also in Ref. [ 211 this fact was explained by different hydrogen incorporation. The presence of a high concentration of diborane in the gas mixture modifies the elemental incorporation, reducing the concentration of hydrogen with respect to that of silicon. Since the implantation process does not modify the elemental composition, the comparison with the Eg values obtained on implanted samples confirms that the enhanced reduction of the Taut gap detected with gas-phase B with respect to P doping is due to differences in film composition rather than to the presence of the boron itself. In Fig. 5 the Urbach energy E, and the deep defect density N,, are reported for the films deposited with 40 seem of CH, for two different B and P concentrations as a function of annealing temperature. The data were obtained from the subbandgap absorption measured by PDS; Nd was computed through the relationship Nd = 1.15 X 10’6a,,z [ 221 where a1,2 is the absorption coefficient at 1.2 eV, to which the absorption value of the exponential tail, extrapolated to the same energy, was subtracted. The results show the expected large overall absorption increase after implantation, and a gradual recovery after annealing, as also indicated by IR measurements. The optimum annealing temperature to obtain a recovery of radiation damage without further degradation due to hydrogen evolution is, for all the films, in the range 250-275 “C. Implantation, after annealing at 250 “C, provides doped samples with an energy gap of about 1.8 eV and defect density of about lo’* cmw3, independent of dopant type, concentration level and carbon content. In a boron-doped sample implanted at 5 X lo*’ cmp3 the broadening of the slope of a(E) is very strong and the Urbach energy remains as high as 200 meV even after annealing. The slope change between tail and defect absorption region is hardly detected, which makes difficult a unique evaluation both of E. and N& This behaviour was already detected for gas-phase doping [ 201 and has therefore to be attributed to
In Fig. 6 the dark conductivity u,+ the conductivity activation energy E, and the pre-exponential factor go are reported as a function of impurity concentration for samples deposited with CH4 flow rates varying from 40 to 70 seem and after annealing at 250 “C. We can observe that, by increasing the doping, the u. values decrease from about 500 to about 20 S cm-‘, indicating the transition from extended state conduction to hopping in the tail states (see for instance Ref. [ 101, pp. 119-120). At concentrations as low as 10” cme3, the samples show a behaviour that is typical of intrinsic samples ( od values lower than 10-i’ S cm-’ and Fermi level at about midgap) whereas, at higher concentrations, the Fermi level shifts towards the bands, and the doping becomes effective. Also, no significant differences are detected in the electrical properties of boron- and phosphorus-implanted samples with the same implanted concentration. We remark that we are comparing films that have the same elemental composition: we can therefore conclude that the doping efficiency for B and P is the same. On the contrary in the case of gas-phase doping [ 201 the same chemical concentration of B and P in the gas mixture results in a two orders of magnitude higher dark conductivity for B-doped samples. However as underlined above (see Section 3.3), in this case the film composition, and consequently the Taut gap, are different and therefore a direct comparison of doping efficiency of B and P cannot be made. The higher conductivity value shown by
Fig. 6. Dark conductivity, activation energy and preexponential factor as a function of dopant concentration for a-SiC:H films deposited with different CH, flow rates, boron and phosphorus implanted, after annealing at 250 “C.
F. Demichelis
et al. /Thin Solid Films 265 (1995) 113-118
117
4. Conclusions
P
B-imp
1.7
1.6 1.5
0
P-imp
0
B-gas
0
P-gas
od
[S cm-‘]
Fig. 7. Taut’s energy gap as a function of dark conductivity for a-SiC:H films deposited with varying CH, flow rate and boron or phosphorus doped by ion implantation at different dopant concentration. In the figure results from gas-phase doped films are also included.
gas-phase B doped samples is attributed to a lower value of the Taut gap. 3.4. Comparison between optical and electrical properties of implanted and gas-phase doped a-SiC:Hfilms The peculiarity of a-SiC:H films is that of being an amorphous semiconductor with good optical properties, in particular with high transparency and conductivity for a variety of optoelectronic applications. For this reason, looking for the best trade-off between high energy gap and high conductivity, in Fig. 7 the energy gap values are reported as a function of dark conductivity at room temperature for B- and P-implanted films after annealing at 250 “C and for PH3 and B21&doped films, with different dopant concentrations and different methane flow rates. It is observable that, within experimental errors, all the implanted films, represented by full symbols in the figure, lie approximately on the same curve. For dark conductivities lower than lo-” S cm-’ the increase in crd is associated with a decrease in the energy gap: ion implantation at concentrations lower than 1Or9cmp3 causes a slight gap shrinkage and a consequent slow increase in conductivity. For higher concentrations, the energy gap stabilizes in the range 1.80-1.90 eV, while the doping effect results in dark conductivity increasing up to 3 X lo-’ S cm-‘. In the same figure values of Eg and gd for samples obtained from gasphase doping are also reported for comparison. For the highest dopant concentration, the dark conductivity values of the implanted samples are the same which have been obtained by gas-phase doping, indicating that the electrical behaviour is mainly governed by the impurity concentration rather than by the incorporation mechanism. However since both chemical and structural composition also govern the Eg - vd plot, a certain spread of the data is obtained, and, for lower conductivities, higher energy gap values are obtained with gasphase doping.
Implantation strongly modifies the structure of a-SiC:H films, promoting monohydride bonds, and changing the microstructure; the damage is only partially recovered by annealing. For all films the best annealing temperature in order to obtain recovery of the radiation damage is in the range 250-275 “C. Implantation causes a gap shrinkage which is quite the same for B- and P-implanted films and for the higher implanted concentrations the energy gap stabilizes at values in the range 1.80-l .90 eV after annealing at 250 “C. This allows one to conclude that the enhanced gap shrinkage normally detected with gas-phase boron with respect to phosphorus doping is to be attributed to the different elemental incorporation during deposition. Sub-bandgap absorption of implanted films after annealing is similar to that detected for gas-phase doping, indicating that it is mainly affected by dopant concentration rather than by residual radiation damage. For all implanted films the defect density as measured from PDS stabilizes at about 10’s cme3 after annealing at 250 “C and an increase of the Urbach energy with the implanted concentration, stronger for boron implantation, is observable. Boron and phosphorus appear to have the same doping efficiency in implanted samples. This seems to be in contrast with the results obtained with gas phase-doping [ 201. However in the case of gas-phasedoping, the higher conductivity value of B-doped films is shown to be a consequence of a lower value of the Taut gap.
References [ I] Y. Tawada, K. Tsuge, M. Konde, H. Okamoto and Y. Hamakawa. J. Appl. Phys. 53 (1982) 5273. [2] Y. Li, in M.J. Thompson, A. Madan, K. Tanaka, P.G. LeComber and E. Schiff (eds.), Amorphous Silicon Technology, Vol. 297, Material
Research Society, Pittsburgh, PA, 1992, p. 803. [3] J. Kanicki, Amorphous and Microcrystalline Semiconductor Devices, Artech House, Norwood, 1991. [4] N. Basu, G. Ganguly, S. Ray and A.K. Barna, Jpn. J. Appl. Phys., 28 (1992) 1776. [5] Y. Tawada, M. Kondo, H. Okamoto and Y. Hamakawa, Proc. 15th IEEE Photovoltaic Specialist Conference, FL, IEEE, New York, 1981, p. 245. [6] F. Fujimoto, A. Ootuca, K. Komaki, Y. Yaata, I. Ymane, H. Yamashita, Y. Hashimoto, Y. Tawada, K. Nishimura, H. Okamoto and Y. Hamakawa, Jpn. J. Appl. Phys., 23 (1984) 810. [7] Y. Kuwano and S. Tsuda, Amorphous and Microcrystalline Silicon Carbide, Vol. 34, Springer Proc. in Physics, Springer-Verlag. Berlin, 1989, p. 167. 181 S. Kalbitzer, G. Muller, P.G. LeComberand W.E. Spear, Philos. Mag. B, 41 (1980) 439. [9] S. Vavilov, I.P. Akimchenko and V.V. Krasnopevtsev, Phys. Status Solidi (a), 94 ( 1986) 343. [ 101 A. Madan and M.P. Shaw, The PhysicsandApplicationsofAmorphous Semiconductors, Academic Press, New York, 1988.
F. Demichelis et al. / Thin Solid Films 265 (1995) I1 3-1 I8
118
[ 111 G. Muller and PG. LeComber, Phifos. Mag. B, 39 ( 1981) 159. [12] R. Rizzoli, R. Galloni, C. Summonte, F. Demichelis, C.F. Pini, E. Tresso, G. Crovini, P. Rava and F. Zignani, Amorphous Silicon Technology, Material Research Society, Pittsburgh, PA, 1994, in press. [13] R.Galloni,Y.S.Tsuo,D.W. BakerandF. Zignani,J. Non-Cryst.Solids, 114 (1989) 271. [ 141 R. Gal1oni.M. Ruth, A. Desalvoand Y.S. Tsuo, PhysicaE, 170 (1991) 273. [ 151 F.J. Fonseca, R. Galloni and A. Nylandsted Larsen, Philos. Mag. B, 67 (1993) 107. [ 161 F. Demichelis, R. Galloni, C.F. Pi,
R. Rizzoli, C. Summonte and E. Ttesso, in J. Thompson, Y. Hamakawa, P.G. LeComber, A. Madan and E. Schiff (eds.), Amorphous Silicon Technology, Materials Research Society, Vol. 258.1992, p. 93.
[ 171 R. Rizzoli, R. Galloni, C. Summonte, F. Demichelis, C.F. Pi, E. Tresso, G. Crovini, P. Rava and F. Zignani, in E.A. Schiff, M. Hack, A. Madan, M.J. Powell and A. Matsuda (ed.), Amorphous Silicon Technology, Materials Research Society, in press. [ 181 J. Bullot and M.P. Schmidt, Phys. Status Solidi (b), 143 (1987) 345. [ 191 F. Demichelis, C.F. Pirri, E. Tmsso, G. Della Mea, V. Rigato, P. Rava and G. Amato, in A. Madan, M.J. Thompson, P.C. Taylor and P.G. LeComber (eds.),AmotphousSilicon Technology. Vol. 219, Materials Research Society, Pittsburgh, PA, 1991, p. 793. [20] F. Demichelis, C.F. Pirri, E. Tresso, G. Della Mea, V. Rigato and P. Rava, Semicond. Sci. Technol., 6 (1991) 1141. [21] CC. Tsai, Phys. Rev. B, 19 (1979) 2041. [22] N. Wyrsch, F. Finger, T.J. McMahon and M. Vanecek, J. Non-Cryst. Solids, 137&138 (1991) 347.
Thin Solid Films 265
( 1995) 113-l 18
Boron and phosphorus doping of a-SiC:H thin films by means of ion implantation F. Demichelis a, G. Crovini a, C.F. Pirri a, E. Tress0 a, R. Galloni b, C. Summonte b, R. Rizzoli b, F. Zignani ‘, P. Rava d aDipartimento di Fisica, Politecnico di Torino, Italy b CNR-Lamel, Bologna, Italy ’ Facoltci di Ingegneria. Vniversitri di Bologna, Italy d Elettrorava S.p.a.. Savonera, Torino, Italy Received 21 November
1994; accepted 28 March 1995
Abstract A detailed study of the structural and optoelectronic properties of boron- and phosphorus-implanted a-SiC:H thin films is presented. The films have been deposited by ultra high vacuum plasma-enhanced chemical vapour deposition and have energy gap values of about 2 eV. The effects of varying carbon content, boron-and phosphorus-implanted dose and annealing temperatures are reported. It is found that ion implantation strongly modifies the structure, producing an increase of monohydride bonds. After annealing the defect density is mainly affected by dopant type and concentration rather than by residual radiation damage. We also show that no enhanced gap shrinkage is produced by boron with respect to phosphorus doping if the film composition is the same. A comparison with gas phase doping is also reported. Keywork
Amorphous
materials;
Ion implantation;
Silicon carbide; Phosphorus;
1. Introduction Intrinsic a-SiC:H films are widely used for solar cells, lightemitting diodes, photodiodes, sensors and several other applications [ l-31, taking advantage of the possibility of tailoring their optical gap in the range 1.8-3.0 eV. Thin layers of aSiC:H with high conductivity, high transparency and low defect density are needed for instance in solar cell technology and can be obtained by gas-phase doping with either boron (p-type material) or phosphorus (n-type material) [4,5]. Generally the p-type doping is produced by mixing into the SiH4 gaseous phase B2Hs, while PH3 is used for n-type doping. The properties of a-SiC:H films doped using gases such as B,H, and PH, have been widely investigated [ 6,7]. Post-deposition doping techniques such as ion implantation [ 8-121 offer the advantage of a better control of the impurity concentration and the possibility of creating doped patterns in an undoped matrix. Nevertheless the creation of a great deal of damage due to the implantation process is the major drawback of the ion-implantation technique. Both point defects and voids are originated by the ion bombardment, resulting in a significant increase of the density of localized gap states. A feasible method to recover the damage is to perform an annealing process [ 13-151. 0040~6090/95/$09.50 0 1995 Elsevier Science S.A. All rights reserved SSDIOO40-6090(95)06618-7
Boron
Data about high-dose ion implantation doping on amorphous silicon carbide alloy films with energy gap values in the range 1.94-2.10 eV have already been published [ 16,171. In this paper we present a detailed study on the structural and optoelectronic properties of a-SiC:H thin films deposited by ultra high vacuum plasma-enhanced chemical vapour deposition (UHV-PECVD), implanted either with boron or phosphorus. The evolution of bonding structure, electrical activation and residual damage as a function of carbon content, implanted concentration and annealing temperature are reported.
2. Experimental Samples were deposited by UHV-PECVD in a SiH4 + CH, gas mixture. The deposition conditions are reported in Table 1: they were optimized in order to have device-quality a-SiC:H films with an energy gap of about 2 eV. In order to obtain homogeneous doping through the sample thickness (about 0.3 p.m for all the samples), multiple energy phosphorus- and boron-ion implantation was performed. Doping uniformity better than 5% was obtained by successively implanting at three energies, which were 15, 55, 140
114
F. Demichelis et al. /Thin Solid Films 265 (1995) 113-118
Table 1 Deposition conditions Substrate temperature Total pressure R.f. power CH, flow rate SiH, flow rate
175-190 “C 80 Pa 2.8 lo-’ W cm-* 40-70 seem 40 seem
keV for P ions and 15, 33, 60 keV for B ions. The total implanted peak concentration ranged from lOI to 5 X 10” cmm3. In order to study the recovery of radiation damage after ion implantation, samples were isochronally annealed for 1 h in air at increasing temperatures between 150 and 400 “C. The optical absorption coefficient was obtained by transmittance and reflectance spectroscopy in the fundamental absorption region by means of a Perkin-Elmer UV-Vis-NIR Lambda 9 spectrophotometer in the 0.2-1.5 pm range and by photothermal deflection spectroscopy (PDS) in the subgap region. The nature of chemical bonding was studied by infrared spectroscopy in the absorption mode using a Perkin-Elmer FTIR 2000 spectrophotometer between 400 and 4000 cm- ’ The absorption coefficient was deduced from the transmission spectra taking into account the influence of interference fringes and absorption due to the silicon substrate. The deconvoluted peaks corresponding to the different vibrational modes and their integrated intensities were determined. Steady-state dark conductivity at 27 “C and conductivity versus temperature were measured in a planar configuration in a MMR vacuum chamber equipped with a K20A temperature controller. Different substrates suitable for different characterizations were used: Corning glass 7059 for electrical measurements, optically polished fused silica for optical measurements and PDS and FZ crystalline silicon for infrared (IR) spectroscopy. 3. Resultsand discussion 3.1. Structural properties The IR absorption spectra have been measured on as deposited, as implanted and annealed films with different CH, flow rates. All the spectra show the characteristic vibrational modes typical of a-SiC:H with a low carbon content [ 181 No significant variation as a function of carbon concentration has been observed in the spectra of as-deposited films. In Fig. 1 the deconvoluted absorption spectra, after subtraction of the interference fringes, are reported for the film deposited with 40 seem of Cl& flow rate, and, for the same film, after implantation with boron at different concentrations and after annealing at temperatures between 250 and 350 “C. The arrows indicate the modes of vibration, which are reported in
600
1006'
600
0
2000
2200
[cm-‘]
Fig. 1. Infrared absorption spectra of a-SiC:H films implanted with boron at two different concentrations and annealed at increasing temperature.
Table 2. A similar trend was obtained for the same film when it was implanted with phosphorus. By comparing the spectra of implanted films with those obtained on as-deposited ones it is evident that: there is a broadening of the peak at 640 cm-‘, due to SiH rocking and a decrease of the 780 cm-’ due to the SiCH3 rocking/wagging mode; - there is a disappearance of the doublet structure at 845890 cm-‘, attributed to (SiH,) n scissors modes; there is a decrease of the stretching modes of SiH2 (2070 cm-‘) and a widening of the 2000 cm-’ peak, due to SiH stretching modes. Doping by ion implantation is associated with induced structural changes: in fact the observed trends in the IR spectra depend on the implanted concentrations and are more marked in the heavily implanted films. Ion implantation promotes monohydride bonds at the expense of the initially dominating SiH2 and Six-H vibrations and modifies the characteristic microstructure of the a-SiC:H network. We note that these variations in IR spectra are attributed to radiation damage, as dopant concentrations similar to those used in this paper, when introduced by gas phase doping, do not influence the a-SiC:H IR spectrum [ 191.
Table 2 IR vibrational peak assignments 1 2 3 4 5 6
640 cm-’ 780 cm-’ 845-890cm-’
1000 cm-’ 2010 cm-’ 2070 cm-’
SiH wagging and/or rocking Si-CH3 and/or Sic rocking/wagging SiH, scissors CH, wagging SiH stretching SiH2 stretching
F. Demichelis etal. /Thin Solid Films 265 (1995) 113-118
B imp. 5 10zD cmS3
7 3-
I
Pimp. 1P
\ \
‘2070 .
.\\
.i ‘2wo
._/\/
cm4
3
‘-‘\
2 -’
P imp. 1018 cm-3
daesp. imp. as
250
300
350 400 T
‘annealing
c&
p,“p.
Boron
250
300 350
400
i3
I
[ “‘1
Fig. 2. Integrated absorption intensity of SiH and SiHz stretching peaks for films deposited with 40 seem CH., flow rate and implanted with boron and phosphorus at two different concentrations.
In order to study the radiation damage and its recovery, the integrated intensity, defined as: ff( a) do w
In=]---
has been evaluated for the absorption peaks centred at 20102020 cm- ’ and at 2070 cm- ’ at different annealing temperatures for the film deposited with 40 seem of methane flow and implanted at different boron and phosphorus concentrations. These are reported in Fig. 2. It is evident that the radiation damage is stronger for the higher implanted concentrations, for which the increase of the SiH integrated intensities is so strong to overcome the value of the SiH, one. When the samples are annealed at temperature between 250 and 300 “C a slight decrease of the SiH peak is accompanied by a slight increase of the SiH* one, indicating a tendency towards a structural recovery of the radiation damage. For annealing temperatures higher than 300 “C the decrease of the H-related peaks clearly indicates, in all the films, the onset of hydrogen evolution which, above 400 “C, induces stress and peeling-off of the samples.
13 t
I
I
I
as dep.
as imp.
150
I
200
I
250
I,,
,
300
as as dep. imp.
T annealing
150
.(
200
250
300
[“Cl
Fig. 3. Taut’s energy gaps versus annealing temperature of a-SiC:H films deposited with 40 seem CH, flow rate, implanted with boron and phosphorus at varying concentrations.
maximum value, whereas for higher temperatures the decrease of Eg is due to hydrogen evolution. In Fig. 4, the Eg values are reported for the samples deposited at CH, flow rates of 40 and 70 seem as a function of the annealing temperatures and for two different P and B concentrations. For dopant concentrations as low as lo’* cme3 the energy gap depends on carbon content. For higher implanted concentrations, after annealing at 250 “C, the Eg values lie in the range 1.8-l .9 eV, independent of the CI-L,flow rate. We would like to underline one feature that comes out from our results. As will be shown in the following (see Section 3.3), the doping is electrically effective for concentrations higher than about lOi cmd3. For those concentrations, and after annealing at 250 “C, the gap shrinkage in the case of B or P doping by ion implantation, is approximately the same. In the case of doping by the gas phase, instead, it was previously reported [20] that boron doping is accompanied by a more enhanced gap
3.2. Optical properties In Fig. 3 the optical gap Eg, evaluated from the Taut plot, is reported as a function of annealing temperature for the sample deposited at a CH4 flow rate of 40 seem implanted at four different B and P concentrations. The radiation damage produced by ion implantation results in a gap shrinkage which is more marked for higher implanted concentrations. The amount of damage also depends on the atomic number of the implanted ions: the gap shrinkage detectable, in the case of phosphorus implantation, even at low concentration, is a consequence of the larger weight of P with respect to B ion. The implantation damage is partially recovered by annealing. For each concentration, at annealing temperatures between 250 and 275 “C the Taut gap of implanted films reaches the
:1:1.:_::.1.‘: 1yL~z= I%. Ep. 150
200
250
300
&
Tannealing
#&
150
200
250
300
[ “Cl
Fig. 4. Taut’s energy gaps versus annealing temperature for a-SiC:H films deposited with different CH, flow rates and implanted at two different boron and phosphorus concentrations.
116
F. Demichelis et al. /Thin Solid Films 265 (1995) 113-l 18
the presence of boron atoms rather than to the implantation process. 3.3. Electrical properties
I. 5.
i&.
. 200
. 300
I
400
I. dt;.
,, la,6.
200
, 300
I 400
T annealing [“Cl Fig. 5. Urbach energy and defect densities of films deposited with 40 seem of CH, flow rate and implanted with boron and phosphorus at two different concentrations.
shrinkage than phosphorus doping. In Ref. [ 201 and also in Ref. [ 211 this fact was explained by different hydrogen incorporation. The presence of a high concentration of diborane in the gas mixture modifies the elemental incorporation, reducing the concentration of hydrogen with respect to that of silicon. Since the implantation process does not modify the elemental composition, the comparison with the Eg values obtained on implanted samples confirms that the enhanced reduction of the Taut gap detected with gas-phase B with respect to P doping is due to differences in film composition rather than to the presence of the boron itself. In Fig. 5 the Urbach energy E, and the deep defect density N,, are reported for the films deposited with 40 seem of CH, for two different B and P concentrations as a function of annealing temperature. The data were obtained from the subbandgap absorption measured by PDS; Nd was computed through the relationship Nd = 1.15 X 10’6a,,z [ 221 where a1,2 is the absorption coefficient at 1.2 eV, to which the absorption value of the exponential tail, extrapolated to the same energy, was subtracted. The results show the expected large overall absorption increase after implantation, and a gradual recovery after annealing, as also indicated by IR measurements. The optimum annealing temperature to obtain a recovery of radiation damage without further degradation due to hydrogen evolution is, for all the films, in the range 250-275 “C. Implantation, after annealing at 250 “C, provides doped samples with an energy gap of about 1.8 eV and defect density of about lo’* cmw3, independent of dopant type, concentration level and carbon content. In a boron-doped sample implanted at 5 X lo*’ cmp3 the broadening of the slope of a(E) is very strong and the Urbach energy remains as high as 200 meV even after annealing. The slope change between tail and defect absorption region is hardly detected, which makes difficult a unique evaluation both of E. and N& This behaviour was already detected for gas-phase doping [ 201 and has therefore to be attributed to
In Fig. 6 the dark conductivity u,+ the conductivity activation energy E, and the pre-exponential factor go are reported as a function of impurity concentration for samples deposited with CH4 flow rates varying from 40 to 70 seem and after annealing at 250 “C. We can observe that, by increasing the doping, the u. values decrease from about 500 to about 20 S cm-‘, indicating the transition from extended state conduction to hopping in the tail states (see for instance Ref. [ 101, pp. 119-120). At concentrations as low as 10” cme3, the samples show a behaviour that is typical of intrinsic samples ( od values lower than 10-i’ S cm-’ and Fermi level at about midgap) whereas, at higher concentrations, the Fermi level shifts towards the bands, and the doping becomes effective. Also, no significant differences are detected in the electrical properties of boron- and phosphorus-implanted samples with the same implanted concentration. We remark that we are comparing films that have the same elemental composition: we can therefore conclude that the doping efficiency for B and P is the same. On the contrary in the case of gas-phase doping [ 201 the same chemical concentration of B and P in the gas mixture results in a two orders of magnitude higher dark conductivity for B-doped samples. However as underlined above (see Section 3.3), in this case the film composition, and consequently the Taut gap, are different and therefore a direct comparison of doping efficiency of B and P cannot be made. The higher conductivity value shown by
Fig. 6. Dark conductivity, activation energy and preexponential factor as a function of dopant concentration for a-SiC:H films deposited with different CH, flow rates, boron and phosphorus implanted, after annealing at 250 “C.
F. Demichelis
et al. /Thin Solid Films 265 (1995) 113-118
117
4. Conclusions
P
B-imp
1.7
1.6 1.5
0
P-imp
0
B-gas
0
P-gas
od
[S cm-‘]
Fig. 7. Taut’s energy gap as a function of dark conductivity for a-SiC:H films deposited with varying CH, flow rate and boron or phosphorus doped by ion implantation at different dopant concentration. In the figure results from gas-phase doped films are also included.
gas-phase B doped samples is attributed to a lower value of the Taut gap. 3.4. Comparison between optical and electrical properties of implanted and gas-phase doped a-SiC:Hfilms The peculiarity of a-SiC:H films is that of being an amorphous semiconductor with good optical properties, in particular with high transparency and conductivity for a variety of optoelectronic applications. For this reason, looking for the best trade-off between high energy gap and high conductivity, in Fig. 7 the energy gap values are reported as a function of dark conductivity at room temperature for B- and P-implanted films after annealing at 250 “C and for PH3 and B21&doped films, with different dopant concentrations and different methane flow rates. It is observable that, within experimental errors, all the implanted films, represented by full symbols in the figure, lie approximately on the same curve. For dark conductivities lower than lo-” S cm-’ the increase in crd is associated with a decrease in the energy gap: ion implantation at concentrations lower than 1Or9cmp3 causes a slight gap shrinkage and a consequent slow increase in conductivity. For higher concentrations, the energy gap stabilizes in the range 1.80-1.90 eV, while the doping effect results in dark conductivity increasing up to 3 X lo-’ S cm-‘. In the same figure values of Eg and gd for samples obtained from gasphase doping are also reported for comparison. For the highest dopant concentration, the dark conductivity values of the implanted samples are the same which have been obtained by gas-phase doping, indicating that the electrical behaviour is mainly governed by the impurity concentration rather than by the incorporation mechanism. However since both chemical and structural composition also govern the Eg - vd plot, a certain spread of the data is obtained, and, for lower conductivities, higher energy gap values are obtained with gasphase doping.
Implantation strongly modifies the structure of a-SiC:H films, promoting monohydride bonds, and changing the microstructure; the damage is only partially recovered by annealing. For all films the best annealing temperature in order to obtain recovery of the radiation damage is in the range 250-275 “C. Implantation causes a gap shrinkage which is quite the same for B- and P-implanted films and for the higher implanted concentrations the energy gap stabilizes at values in the range 1.80-l .90 eV after annealing at 250 “C. This allows one to conclude that the enhanced gap shrinkage normally detected with gas-phase boron with respect to phosphorus doping is to be attributed to the different elemental incorporation during deposition. Sub-bandgap absorption of implanted films after annealing is similar to that detected for gas-phase doping, indicating that it is mainly affected by dopant concentration rather than by residual radiation damage. For all implanted films the defect density as measured from PDS stabilizes at about 10’s cme3 after annealing at 250 “C and an increase of the Urbach energy with the implanted concentration, stronger for boron implantation, is observable. Boron and phosphorus appear to have the same doping efficiency in implanted samples. This seems to be in contrast with the results obtained with gas phase-doping [ 201. However in the case of gas-phasedoping, the higher conductivity value of B-doped films is shown to be a consequence of a lower value of the Taut gap.
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Research Society, Pittsburgh, PA, 1992, p. 803. [3] J. Kanicki, Amorphous and Microcrystalline Semiconductor Devices, Artech House, Norwood, 1991. [4] N. Basu, G. Ganguly, S. Ray and A.K. Barna, Jpn. J. Appl. Phys., 28 (1992) 1776. [5] Y. Tawada, M. Kondo, H. Okamoto and Y. Hamakawa, Proc. 15th IEEE Photovoltaic Specialist Conference, FL, IEEE, New York, 1981, p. 245. [6] F. Fujimoto, A. Ootuca, K. Komaki, Y. Yaata, I. Ymane, H. Yamashita, Y. Hashimoto, Y. Tawada, K. Nishimura, H. Okamoto and Y. Hamakawa, Jpn. J. Appl. Phys., 23 (1984) 810. [7] Y. Kuwano and S. Tsuda, Amorphous and Microcrystalline Silicon Carbide, Vol. 34, Springer Proc. in Physics, Springer-Verlag. Berlin, 1989, p. 167. 181 S. Kalbitzer, G. Muller, P.G. LeComberand W.E. Spear, Philos. Mag. B, 41 (1980) 439. [9] S. Vavilov, I.P. Akimchenko and V.V. Krasnopevtsev, Phys. Status Solidi (a), 94 ( 1986) 343. [ 101 A. Madan and M.P. Shaw, The PhysicsandApplicationsofAmorphous Semiconductors, Academic Press, New York, 1988.
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