A national laboratory of the U.S. Department of Energy 

Office of Energy Efficiency & Renewable Energy 

National Renewable Energy Laboratory 

Innovation for Our Energy Future 

Microscopic Mechanism of Subcontract Report 

NREL/SR-520-37990 the Staebler-Wronski Effect in May 2005 

a-Si Films and High-Efficiency 

Solar Cells 

Final Subcontract Report 

1 October 2001—30 September 2004 

D. Han 

University of North Carolina 

Chapel Hill, North Carolina 

NREL is operated by Midwest Research Institute ● Battelle Contract No. DE-AC36-99-GO10337 
Microscopic Mechanism of Subcontract Report 

NREL/SR-520-37990 the Staebler-Wronski Effect in May 2005 

a-Si Films and High-Efficiency 

Solar Cells 

Final Subcontract Report 

1 October 2001—30 September 2004 

D. Han 

University of North Carolina 

Chapel Hill, North Carolina 

NREL Technical Monitor: Bolko von Roedern 

Prepared under Subcontract No. ADJ-1-30630-09 

National Renewable Energy Laboratory 

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PREFACE 

This final Technical Report covers the work performed by University of North Carolina at Chapel Hill 

(UNC-CH) for the period Oct.1, 2001 to Sept. 30, 2004 under subcontract No. ADJ-1-30630-09. The 

report was written by Daxing Han (P.I.) and produced with Keda Wang (co-P.I.) Yue Wu (co-P.I.); 

Jonathan Baugh; Andrea Hilchey, J.R. Weinberg-Wolf, E.C.T. Harley, Chris Lawyer, and Anthony 

Canning, graduate students; and Brittany Huie, under graduate student. The samples were obtained from 

Team members at USOC, BPSolar and the a-Si group at the National Renewable Energy Laboratory 

(NREL). 

This work is part of the amorphous silicon (a-Si:H) teamwork of the National Thin Film PV Program and 

NREL Thin-Film partnership program. Over the last three years, we have benefited from numerous 

discussions and cooperation with our Condensed Matter colleagues at UNC-CH, L. E. McNeil, and Y. 

Wu. At other institutions we have collaborated with: Qi Wang, Yueqing Xu, Lynn Gedvilas, Brent 

Nelson and Paul Stradins et al. at NREL; Jeff Yang and Baojie Yan et al at USOC; and G. Ganguly at 

BPSolar. We have also collaborated with S. Nitta and H. Habuch, Gifu University, and Tatsuo Shimizu, 

Kanazawa University, Japan. Our research has resulted in the publication of 21 papers, 6 papers in press 

and the submission of quarterly and annual reports. We have completed a review article on Photo- and 

electro-luminescence of a-Si:H and mixed-phase alloys for a special issue of “Solar Energy Materials and 

Solar Cells” (Elsevier Science) being edited by Hellmut Fritzsche and Subhendu Guha. This article 

summarizes our recent 10-year work of electoluminescence, the dispersive-transport-controlled 

recombination, and photoluminescence in a-Si:H and mixed-phase alloys. 

i 

TABLE OF CONTENTS

Preface i 


Table of Contents ii


List of figures v 

List of tables vii 

Results 

I. Nano-voids Related Phenomena in High Growth Rate Intrinsic Hydrogenated Amorphous 

Silicon Films Prepared by Hot-wire CVD film 1 

I.1 Large red shift of PL peak energy in high-growth rate a-Si:H film
 1 

1.1.1 Samples and experimental
 1 

1.1.2 Results and discussion 4 


1.1.3 Summary
 4 

I.2 Confinement effect on dipole-dipole interactions in nanogas system- NMR results 
 4 

1.2.1 Introduction 4


1.2.2 Sample and experimental 
 4 

1.2.3 Results and discussion 5


I.3 Conclusions 6 


II. Microcrystalline transition from a-Si as a function of hydrogen dilution and substrate

temperature of hot-wire CVD 7 

II.1 Introduction 7


II.2 Samples and Experimental Conditions
 7 

II.3 Results and Discussion 8 


2.3.1 Raman Spectra 8 


 2.3.2 IR Absorption
 9 

2.3.3 Optical Absorption Spectra 13 


2.3.4 Conductivity Activation Energy 15 


2.3.5 PL spectra at 80 K 16 


II.4 Summary
 16 

II.5. Conclusions
 17 

III. Do

 May Not Be The Dominant Recombination Centers For Photoconductivity In Hot-Wire

a-Si:H 17

III.1 Introduction 17


III.2 Sample and Experimental
 19 

III.3 Results 20


3.3.1 Natural degradation effect
 20 

3.3.2. Strong recombination center in a-Si:H prepared by HW-CVD 21


3.3.3. ESR and CPM studies 22 


III.4 Summary and Discussion 23 

ii 

3. 4.1 Summary
 23 

3. 4.2 Natural degradation effect
 23 

3. 4.3 Why the activation energy is very large and driven down?
 24 

III.5 Conclusions 24 


IV. PL study of growth rate dependence on solar cells made using DC-CVD by BPSolar 25


IV.1. A special microstructure of high growth rate films 25 


4.1.1. A special case of the IR spectra 25 


4.1.2. A special case of the PL spectroscopy 27 


4.1.3. Summary of the special case 28 


IV.2 PL results from a group of new samples 28 


4.2.1. Sample and Experimental
 28 

4.2.2. PL Results 29 


4.2.3. In-situ light soaking
 30 

IV.3. Summary and Discussion 
 30 

V. Systematical Raman study of thin films from amorphous-to-nanocrystalline silicon

prepared by HW-CVD 31 

V.1. Introduction 31 


V.2. Sample and Experimental
 32 

V.3. Results and discussions 33


 V.4. Summary
 37 

V.5 Conclusion 38


VI. Micro-Raman Studies of Mixed-Phase Hydrogenated Silicon Solar Cells 

VI.1. Introduction 38 


VI.2. Sample and Experimental
 39 

VI.3. Results and discussions 39


VI.4. Summary
 41 

VI.5. Conclusion 42


VII. Toward understanding the degradation without light soaking in hot-wire a-Si:H thin films 

and solar cells 42


VII.1 Introduction 42


VII.2 Experimental Details
 42 

VII.3 Result and discussion 43


VII.4 Conclusions 46


VIII. Micro-Raman Studies for Electrical-induced Structural Changes in Metal/μ c-Si/Metal

Structures Made by Hot-Wire CVD 46 

VIII.1. Sample and Experiment 
 46 

VIII.2. Results 47


8.2.1. The I-V characteristics of M/a-Si/M and M/μc-Si/M structures
 47 

8.2.2. Raman studies before and after switching of the M/μc-Si/M structures
 48 

VIII.3. Summary 
 48 

iii 

VIII.4. Conclusion 49


IX. Buffer-layer Effect on Mixed-Phase Cells Studied by Raman and PL Spectroscopy 49 

IX.1. Introduction 49


IX.2. Sample and Experimental
 49 

IX.3. Raman results 50


IX.4. PL results 51


9.4.1. PL spectra at low temperature 
 52 

9.4.2. PL intensity temperature dependence 
 52 

9.4.3. in-situ light soaking effect
 53 

IX. 5. Summary 
 53 

IX. 6. Conclusion 54


X. Correlation of Hydrogenated Nanocrystalline Silicon Microstructure and Solar Cell 

Performance 54


X.1. Introduction 54


X.2. Sample
 54 

X.3. Raman and PL results 55


X.4. Discussion 57 


X.5. Summary
 58 

X.6. Conclusion 58


FURTHER EFFORTS 59


Publications 60 


References 62


ABSTRACT 63 


iv

List of figures 

Fig. 1.1	(a) Normalized PL spectra excited by a 632 nm laser, (b) the PL peak energies excited by both a 

632 nm (solid circles) and a 488 nm laser (open circles) at 80 K as a function of the deposition 

rate for HW a-Si:H films with growth rates from 10 to 115 Å/s. 

Fig. 1.2 PL spectra as a function of temperature for two HW films with growth rates 10 Å/s. 

Fig. 1.3 The PL main band and defect band from a typical high growth rate film L356. (a) The integrated 

intensity and (b) the peak energy positions plotted as a function of temperature. 

Fig. 1.4 Raman shift for a-Si:H samples (a) on c-Si substrates, and (b) on ss substrates. 

Fig. 1.5. Illustration of a cross-section of a-Si:H thin-film with a magnified region showing an elongated 

nanovoid containing a high-density H2 gas (blue dots). Such voids show strong alignment of the 

long axis in the growth direction (along the film normal) in the high-growth rate films studied 

here. The angle Ω is therefore the angle between the external magnetic field (H0) and the 

principal axis a of the nanovoid. 

Fig. 1.6. Room temperature proton NMR spectra (a) and their line fits (b) for an a-Si:H thin-film taken at 

different film orientations with respect to the applied magnetic field (4.7 Tesla). The angle Ω is 

defined in Fig. 1.5. 

Fig.1.7 Measured dipolar broadening of the o-H2 peak due to gas-phase intermolecular dipolar 

interactions, versus Ω. The solid line shows the function 4 60Hz P (cos Ω) . 2 

Fig.2.1 Raman shifts from group I and II films with (a) varied hydrogen dilution ratio, R, and (b) varied 

substrate temperature Ts. The Raman spectrum from the reference film is shown on the bottom 

of Fig. 2.1(a). 

Fig. 2.2 The IR spectra in three frequency ranges of 550-750 cm-1, 750-1200 cm-1and 1800-2200 cm-1 that 

cover the Si-H wag mode, the SiH2 bending mode, Si-O modes, and the stretch modes. for (1a, 

1b, 1c) group I and (2a, 2b, 2c) group II films. The values of R and Ts are indicated in the 

legend. 

Fig. 2.3 The correlation of CH (�) and the relative intensity of the 2090 cm-1 peak I2090/(I2000+I2090) (�) 

with microcrystallinity for (a) group I and (b) group II films. 

Fig. 2.4 Optical absorption spectra obtained by using CPM and fitted by TRS for (a) group I and (b) 

group II films. The values of R and Ts are indicated in the legend. The absorption spectrum for 

c-Si is plotted in (a) for reference. 

Fig. 2.5 the optical absorption coefficient at 1.4 eV (�) and the c-Si volume fraction (∆) as a function of 

(a) R, and (b) Ts. 

Fig. 2.6 Position of Fermi-level (�) and the c-Si volume fraction (∆) as a function of (a) R, and (b) Ts. 

Fig. 2.7 Relative PL spectral intensity at 80 K for (a) group I and (b) group II films. The PL spectrum for 

the reference film is show on the bottom of Fig. 1(a). 

Fig. 3.1 	Distribution function (a) for strong recombination and (b) for weak recombination relative to retrapping. Ed is the thermalization demarcation energy, Eq is the quasi-Fermi level and Ec is the 

electron mobility edge. 

Fig. 3.2 	The conductivities before, during and after light soaking for typical intrinsic a-Si:H films made 

by HW-CVD (solid line) and PE-CVD (dotted line). 

Fig. 3.3	Temperature dependence of photo- and dark conductivity before and after 200 mW/cm2 white 

light soaking for 4 h, (a) for HW-CVD film T779 and (b) for PE-CVD film A7119. 

Fig. 3.4 Thermostimulated conductivity for the same samples in Fig. 3 at annealed State A and lightsoaked State B, (a) for HW-CVD film T779 and (b) for PE-CVD film A7119. 

Fig. 3.5 	The density of defects for HW- and PE-CVD films before and after 200 mW/cm2 white light 

soaking for 30 h. (a) Measured by CPM, (b) measured by ESR, and (c) light-induced metastable 

defects measured by CPM and ESR. 

Fig. 4.1 Raman shift for samples with growth rates of 1, 3 and 10 Å/s. 

v 
Fig. 4.2 IR spectra for samples of (a) L0172-3, L0172-4, and L0173-1, (b) L0276-1c, L0277-1c, and 

L0278-1c. 

Fig. 4.3 PL results for samples L0276-1, L0277-1, and L0278-1 (a) spectra at 80 K, (b) spectra at 180 K. 

Fig. 4.4 PL peak energy as a function of temperature for samples L0276-1, L0277-1, and L0278-1. 

Fig. 4.7 Peakmain of samples L2127, L2129, L2316 and L2317 as a function of temperature. 

Fig. 4.8 Imain of samples L2127, L2129, L2316 and L2317 as a function of temperature. 

Fig. 4.9 Relative intensity of the Idefect as a function of temperature for the four samples. 

Fig. 4.10 Relative intensity of the Idefect at 260 K as a function of the sample growth rate before and after 

in-situ light soaking. 

Fig. 4.11 the Peakmain, Imain, and Idefect vs. T for sample L2127 at State A and State B. 

Fig. 5.1. Xc as a function of R at SiH4 flow rates of 3, 8, 16 and 22 sccm. The solid and dotted lines 

indicate the Xc was obtained from the top or bottom layer of ≥ 60 nm, respectively. 

Fig. 5.2. Raman shift and their fittings from the films with SiH4 flow of 3 sccm, TVP=45%, (a) R=1, and 

(b) R=2. And the enlarged TO mode of the data in 2(c) and 2(d). 

Figs. 5.3.(a) FWHM and (b) ωa-Si of the TO mode as function of R. The solid dots, dashes, open circles, 

and crosses represent the films with a SiH4 flow rate of 3, 8, 16, and 22 sccm, respectively. 

There are two regimes separated at R≈3: the fast changing regime when R <3 and the saturated 

regime when R>3. 

Fig. 5.4. ωc-Si as a function of R. 

Fig. 5.5. ωg.b. as a function of R. 

Figure 6.1. Schematics of (a) 4 scans for all samples, and (b) 20 scans for 13678-42. The diameter of the 

dot is ~3 μm. 

Fig. 6.2 Raman spectrum of 13687-32. 

Fig. 6.3 Raman spectrum of 13678-13. 

Fig. 6.4. Raman shifts of mixed-phase cells with (a) VOC = 0.767 V and (b) VOC = 0.748 V. 

Fig. 6.5. Raman Shift of 20 scans on a mixed phase cell 13678-42. 

Fig. 7.1 The hot-probe data for sample T837. 

Fig. 7.2. (a) Current-temperature dependence and (b) current as a function of annealing time at 420 K, for 

samples of H1441 and H1442 with- and without SiNx passivition layer. The samples were 

mounted in the vacuum and in pure oxygen or nitrogen, respectively. 

Fig. 7.3 Degradation experiments: (a) the dark current at 300 K for samples of H1540a in vacuum (×), 

H1540b in pure O2 (�), and H1540c in the air (∆). The data following the arrows indicate the 

current after Ea measurements. (b) Several measurements of Ea obtained from those three 

samples after the current measurements at 300 K for about a month. The data following the 

arrow indicate the measurements were done in vacuum for H1540c. 

Fig. 7.4 PC temperature dependence for (a) HW film T779 and (b) PE-CVD film A7119. 

Fig. 8.1. Typical switching voltage vs. current traces in (a) a 80 μm-diameter Cr/p+

a-Si:H/Ag and (b) a ~ 

5×5 μm2 Cr/p+ μc-Si/Ag devices. 

Fig. 8.2. The Raman spectra fitted using three functions for (a) the seven non-switched, (b) the eight 

switched devices, and (c) comparing the spectra in (a) and (b). 

Fig. 9.1. Raman spectra for the cells on (a) L13869, (b) L13868, and (c) L13870, corresponding to no 

buffer layer, 100 Å thick a-Si:H buffer layer, and 500 Å thick a-Si:H buffer layer, respectively. 

Fig. 9.2. PL spectra at 80 K for cells 11, 13, and 22 on sample 13869 for (a) linear and (b) semi-log plots. 

Fig. 9.3. PL spectra of the low Voc cells of L13869-11, L13868-11, and L13870-11 for (a) linear and 

(b) semi-log plots. 

Fig. 9.4. PL intensities vs. temperature for cells of (a) L13869, (b) L13868, and (c) L13870, where circles 

represent the a-Si:H main band, triangles the g.b. band, and squares the defect band. 

Fig. 9.5 PL spectra at 80 K before and after light soaking for cells of (a) L13869-11, (b) L13868-11, and 

(c) L13870-11. 

Fig. 10.1. Raman spectra of the eight nc-Si:H solar cells. 

vi

Fig. 10.2. PL spectra of the same nc-Si:H solar cells at 80 K excited by 632.8-nm (solid) and 442-nm 

(dashed) laser beams. 

List of tables 

Table I.1. Characteristics of a-Si:H samples with high growth rate.


Table II.1. Preparation conditions and parameters of the transition films prepared by HW-CVD.


Table II.2 The energy peak position and the bandwidth of the Si-H vibration modes.


Table III.1. Sample preparation conditions.


Table III.2. Conductivity activation energies at initial- annealed- and light soaked states.


Table IV.1. Sample preparation conditions, i-layer properties and their solar cell performances.


Table IV.2. Hydrogen content and the absorption strengths of Si-H vibration modes.


Table IV.3 The sample preparing conditions and their PL parameters.


Table V1.1 Sample VOC with Raman results 


Table VI1.1. Sample information 


Table IX.1. The initial performance of the mixed-phase solar cells with different a-Si:H buffer layers at


the i/p interface.


Table X.I. Preparation conditions and solar cell performance of the nc-Si:H solar cells studied.


Table X.2. Summary of the Raman and PL analyses of the nc-Si:H solar cells.


vii

I. Nano-voids Related Phenomena in High Growth Rate Intrinsic Hydrogenated Amorphous 

Silicon Films Prepared by Hot-wire CVD 

1.1. Large red shift of PL peak energy in high-growth rate a-Si:H films[1] 

1.1.1 Sample and experimental: The growth rates and the characteristics of the hydrogenated amorphous 

silicon (a-Si:H) films are listed in Table I. As a reference, a standard device-quality hot-wire chemical 

vapor deposition (HW-CVD) film was deposited on quartz substrate. For the following 

photoluminescence (PL) measurements, either a 632 or a 488 nm wavelength laser beam with a lightchopper at frequency 17 Hz was used. The sample was mounted on a cold stage. The steady state PL 

spectra were analyzed using a grating monochromator equipped with a liquid nitrogen-cooled Ge 

detector. The signals were recorded by a lock-in amplifier. 

Table I.1. Characteristics of a-Si:H samples with high growth rate. 

Sample Growth Mass density Eopt PLpeak at 1

H-NMR Nanovoid Raman TO mode 

ID rate (% of c-Si (eV) 80 K very narrow line volume fractiona Peak FWHM 

(Å/s) density) (eV) (%) (%) (cm-1) (cm-1) 

Standard 

L388 

L256 

L354 

L353 

L252 

L355 

L259 

L356 

10 

30 

50 

55 

80 

90 

100 

110 

115 

98.0 

96.0 

95.8 

95.3 

95.1 

1.68 

1.65 

1.67 

1.65 

1.66 

1.36 

1.23 

1.05 

1.10 

1.14 

1.12 

1.14 

1.14 

1.14 

Absence 

36 

27 

24 

20 

<0.01 

1-2 

~2 

~2 

~2 

2-3 

~2 

2-3 

~2 

480 

" 

" 

" 

" 

" 

" 

" 

" 

70 

" 

" 

" 

" 

" 

" 

" 

" a From both mass density and reference 

1.1.2 Results and discussion: 

0.6 0.8 1.0 1.2 1.4 1.6 

Nor. PL intensity (a.u.) 

Photon Energy (eV) 

(a) 

50 Å/s 55-115 30 10 Å/s

1.0 

1.1 

1.2 

1.3 

1.4 

0 

632 nm laser 

488 nm laser 

E peak (eV) 

Growth rate (Å/s) 

(b) 

/s. 

20 40 60 80 100 120 

Fig. 1.1 (a) Normalized PL spectra excited by a 632 nm laser, (b) the PL peak energies excited by both a 

632 nm (solid circles) and a 488 nm laser (open circles) at 80 K as a function of the deposition rate for 

HW a-Si:H films with growth rates from 10 to 115 Å

1 
In typical intrinsic a-Si:H made by plasma-enhanced (PE)-CVD, the PL spectrum at 80 K is dominated by 

the tail-to-tail radiative transitions that give a PL main-band peaked at ~1.4 eV with full width at half 

maximum (FWHM) ≈ 0.3 eV. A defect-band peaked at ~0.9 eV is 3-4 orders of magnitude weaker than 

the main band at low temperatures. As the temperature is increased, the PL intensity (Ipl) is quenched by 

non-radiative recombination, and the main-band peak energy position (PLpeak) decreases due to carrier 

thermalization down to the lower-energy band tail states. The PL features of the standard low-deposition 

rate HW-CVD film are similar to those of PE-CVD film. Surprisingly, the main peak energy PLpeak is 

extremely low at 80 K in the high-growth rate films compared to the standard film. Fig. 1.1a shows 

normalized PL spectra excited by a 632 nm laser. The PLpeak decreases rapidly from 1.36 to 1.05 eV when 

the growth rate increases from 10 to 50 Å/s; it then reaches a saturated value of ~ 1.1 eV for growth rates 

> 50 Å/s. The PL red shift is as large as 0.25 to 0.30 eV. One may question whether the excitation photon 

energy of the 632 nm laser is large enough for achieving intrinsic excitation. Therefore, we also used a 

488 nm laser to excite the PL. Fig. 1.1b plots the PLpeak excited by both 632 and 488 nm laser at 80 K as a 

function of the deposition rate. Using the 488 nm laser, we observed a PLpeak position even 10-30 meV 

lower than that obtained using the 632 nm excitation. This can be understood as resulting from film nonuniformity in the growth direction and the more shallow penetration of the 488 nm laser beam probing a 

film portion close to the surface. The PL efficiency, on the other hand, is similar to that of the standard 

film in all the high-growth rate films. 

To explore whether or not the PL follows carrier 

115 Å/s 10 Å/s thermalization in band tails, we measured the PL 

80 K 

120 K 

160 K 

200 K 

220 K 

240 K 

260 K 

300 K 

Nor. PL intensity (a.u.) 

0.6 0.8 1.0 1.2 1.4 1.6 

Photon Energy (eV) 

Fig. 1.2 PL spectra as a function of temperature 

for two HW films with growth rates 10 Å/s. 

temperature dependence. Fig. 1.2 shows normalized 

PL spectra excited by a 488 nm laser as a function of 

temperature for the film deposited at 115 Å/s. The 

standard film PL spectra are plotted as a reference. 

At low temperatures, the PL main-band dominates. 

We see clearly that the PLpeak red shifts. As the 

temperature increases, the 0.8-eV defect band 

gradually dominates, and the PL spectral lineshapes 

of the two films become similar. The data was fitted 

using two Gaussian functions f1 and f2 with line 

widths of 0.20 eV and 0.25 eV representing the main 

and defect bands, respectively. Figure 1.3a shows the 

thermal quenching of the PL main-band intensity and 

a weak temperature dependence of the defect band 

intensity. Figure 1.3b shows that the PLpeak1 of the 

PL main band decreases from 1.15 to 1.04 eV as the 

temperature increases from 80 K to 180 K and 

saturates at ~1.04 eV (at which the carriers reach the 

bottom of the tail states). On the other hand, the peak 

energy position of the defect band, PLpeak2, decreases 

from 0.90 to 0.82 eV as the temperature increases 

from 150 to 300 K. The results in Fig. 1.3 indicate 

that both the PL intensity and the peak energy follow 

the model of carrier thermalization in band tail 

states. However, the large red shift of the PL mainband (from tail-to-tail transition) is puzzling since 

the optical gap of the high rate film is similar to that 

of the standard film. 

2 

3 

To explore how the microstructure could be related to the PL red shift, we first measured the Raman shift. 

The spectra in Figs. 1.4(a) and 1.4(b) show the same lineshapes as standard device-quality a-Si:H, i.e. 

transverse optical (TO) mode peaked at ~480 cm-1 with FWHM ~70 cm-1 for all the films in Table I. We 

further studied the hydrogen microstructure and the material structural ordering by 1

H-nuclar magnetic 

resonance (NMR) and mass density measurement. We found a very narrow Lorentzian component in the 

spectra with ~1 kHz linewidth in addition to the conventional Lorentzian (~3 kHz) and Gaussian (~40 

kHz) components in the high growth rate films L353-L356. This very narrow component is related to 

molecular hydrogen that is absent in the standard film but comprises 20-30% of the total spectral intensity 

for the four samples L353-356. The mass densities were obtained by the flotation method. The average 

10-4

10-3

10-2

10-1

100

101

102

50 100 150 200 250 300 350

Intl 1

Intl 2

Integrated PL intensity, IPL

T (K)

(a)

0.80

0.85

0.90

0.95

1.00

1.05

1.10

1.15

50 100 150 200 250 300 350

Peak1

Peak2

Epeak (eV)

T (K)

(b)

Fig. 1.3 The PL main band and defect band from a typical high growth rate film L356. (a) The 

integrated intensity and (b) the peak energy positions plotted as a function of temperature. 

250 300 350 400 450 500 550 600 650

L354 

L353 

L355 

L356 

Raman Shift (cm-1)

(a)

70 cm-1

 

250 300 350 400 450 500 550 600 650

259.nml

252.nml

256.mnl

Raman Shift (cm-1)

(b)

70 cm-1

Fig. 1.4 Raman shift for a-Si:H samples (a) on c-Si substrates, and (b) on ss substrates. 
was 2.28 ± 0.02 and 2.22 ± 0.02 (g/cm3

) for the standard films and the high growth rate films, 

respectively. We obtain a ~2 % volume fraction of the nanovoids in the high growth rate films. This H2 

peak exhibits a long room-temperature spin-lattice relaxation time (T1) ~ 1.2 s, which implies that there is 

a negligible density of spin-lattice relaxation centers (i.e. dangling bonds) associated with the nanovoids. 

This lack of dangling bonds may indicate an ease of weak bond formation on the surfaces. Additionally, 

the spin-spin relaxation time (T2) of the H2 peak shows a dependence on the film orientation with respect 

to the dc magnetic field. Such unusual anisotropy in a gas system can only be understood if we consider 

the effects of confinement in nano-scale voids that are non-spherical and on the average aligned in the 

film. The experimental NMR results of the H2 peak fit well with the calculated properties of H2 in an 

elongated elliptical geometry of 1-3 nm diameter and length several times the diameter. Meanwhile, we 

have learned that small-angle x-ray scattering (SAXS) also indicates that the 1-3% volume fraction 

nanovoids are elongated and aligned in the high growth rate films. 

1.1.3 Summary: A large red shift of the PL peak energy is found in a-Si:H films prepared by HW-CVD 

with a high growth rate ≥ 50 Å/s. The origin of the red shift is clarified by employing 1

H nuclear magnetic 

resonance (NMR) and mass density measurements. A ~2% volume fraction of tube-like nanoscale voids 

is identified. We suggest that highly strained bonds on these surfaces form broad conduction-band tail 

states that are responsible for the PL red shift. 

I. 2. Confinement Effect on Dipole-dipole Interactions in Nanogas System -NMR Results[2] 

I.2.1 Introduction: Intermolecular dipole-dipole interactions were once thought to average to zero in gases 

and liquids as a result of rapid molecular motion, which leads to sharp nuclear magnetic resonance 

(NMR) lines. Recent papers have shown that small residual couplings survive the motional averaging if 

the magnetization is nonuniform or nonspherical. Here, we show that a much larger, qualitatively 

different intermolecular dipolar interaction remains in nanogases and nanoliquids as an effect of 

confinement. The resulting dipolar coupling is identical for all spin pairs and depends on the volume of 

the gas/liquid container, its shape and orientation with respect to the magnetic field. This nanoscale effect 

is quite useful in the determination of nanostructures and could have unique applications in the 

exploration of quantum space. 

I.2.2 Sample and experimental: Proton NMR is carried out on a pulsed NMR spectrometer operating at a 

radio frequency of 200 MHz in an external magnetic field of 4.7 Tesla. The low proton background was 

achieved by employing homemade plastic-free capacitor in the resonant tank circuit of the probe. Only 

carefully etched and cleaned metal or metal-shielded parts are placed in the region of homogeneous 

magnetic field. The sample region was purged by constant flow of dry nitrogen gas to avoid moisture in 

the air. For the measurement of a single a-Si:H film on quartz substrate a flattened solenoid NMR coil 

was employed. The flat coil surface orientation can be varied with respect to the external magnetic field 

for the measurement of orientation dependence. The typical employed 90° rf pulse is 1.5 μs. The 

temperature variation was achieved by placing the rf coil in a resistive furnace. In order to remove the 

effect of the heating current which could generate magnetic field, the current power supply was controlled 

by the pulse programmer of the spectrometer through a TTL pulse, which turns on or off the heating 

current in about 50 ms. The heating current was off during the NMR data acquisition time of about 16 ms 

and was turned back on immediately thereafter. Stable and homogeneous temperature across the sample 

can be achieved by this method over a long time period. 

4 

Fig. 1.5. Illustration of a cross-section of a-Si:H thin-film with a magnified region showing an elongated 

nanovoid containing a high-density H2 gas (blue dots). Such voids show strong alignment of the long axis 

in the growth direction (along the film normal) in the high-growth rate films studied here. The angle Ω is 

therefore the angle between the external magnetic field (H0) and the principal axis a of the nanovoid. 

I.2.3 Results and discussion: 

4 2 0 -2 -4

Intensity (a.u.) 

kHz 

0.0 o 

54.7o 

77.2o 

(a) 

4 2 0 -2 -4

Intensity (a.u.) 

kHz 

0.0 o 

54.7o 

77.2o 

(b) 

Ω

Fig. 1.6. Room temperature proton NMR spectra (a) and their line fits (b) for an a-Si:H thin-film 

taken at different film orientations with respect to the applied magnetic field (4.7 Tesla). The 

angle is defined in Fig. 1.5. 

It is well known that a-Si:H possesses nanovoids, which contain high-pressure H2 gas of up to 2 kbar. 

Recently, it was found [2] that such nanovoids are elongated and aligned perpendicular to the film surface 

in a-Si:H produced by HW-CVD at high growth rate, as illustrated in Fig. 5. Small angle x-ray scattering 

5 

(SAXS) shows that the total volume fraction of such nanovoids is ~2% and that the estimated diameters 

of the nanovoids are a few nm and the lengths are several times longer than the diameter. To investigate 

the angular dependence NMR measurements were carried out on a single a-Si:H film about 1 μm thick on 

quartz substrate using a flattened solenoid coil in an ultralow 1

H background probe. Figure 6 shows the 

room temperature (RT) 1H spectra of high growth rate (55 Å/s) HW-CVD a-Si:H at Ω=0°, 55°, and 77°. 

The inset of Fig. 6 shows both the conventional broad (FWHH=35 kHz) and narrow (FWHH=3.3 kHz) 

peaks associated with Si-H clusters and isolated Si-H bonds, respectively. In addition, a narrower peak 

with FWHH of less than 1 kHz is also visible. This peak arises from H2 in nanovoids and broadens 

significantly below 60 K due to the slow-down of H2 motion and subsequent H2 freezing. Fittings of the 

conventional narrow peak and the H2 peak are also shown in Fig. 1.6. 

The contribution of the dipole-dipole interactions to this linewidth was determined by the Hahn-echo 

decay using the pulse sequence 0o9 − t − o180 − t − echo . As expected, the rate of the echo height decay 

is the slowest at Ω=55° and follows exp( −2 t / T2 ) with T2=1.3 ms. The Ω−dependent part of the Hahnecho decay was determined from the Gaussian part of the decay function . Figure 7 shows the 

measured 36.2 M 2 / 2π versus the angle Ω. The experimental data are consistent with the calculation 

very well. 

0 

100 

200 

300 

400 

500 

600 

0 

Hz 

Ω (degrees)

20 40 60 80 

Fig.1.7 Measured dipolar broadening of the 

o-H2 peak due to gas-phase intermolecular 

dipolar interactions, versus Ω. The solid line 

shows the function 4 60Hz P (cos Ω) . 2 

The NMR results clarified that the nanovoids are not only related to dangling bond defect centers but 

could be related to radiative centers. We propose that highly strained bonds (weak bonds) on the surfaces 

of the tube-like nanovoids result in a broad conduction band tail. In a recent molecular dynamics study of 

band tails in a-Si:H, Fedders et al [3] conclude that stretched bonds affect the conduction band tail; 

whereas, short bonds affect valence band tail states. The nanovoid-related states do not produce a new 

type of luminescence center but only contribute to the continuous distribution of tail states. Therefore, 

such states could cause the peak energy to red shift instead of forming a new PL peak. [1]. 

I.3 Conclusions: 

We observed new phenomena in high growth rate HWCVD a-Si:H films by PL and NMR. We 

explain that the nonsocial elongated nanovoids filled with one-dimensional H2 gas alignment in the 

growth direction in the high-growth rate films. 

6 

II. Microcrystalline transition from a-Si as a function of hydrogen dilution and substrate 

temperature of hot-wire CVD [4] 

II.1 Introduction 

The transition films from amorphous (a)- to microcrystalline (μc) -Si:H have attracted a great deal of 

attention because the most stable, high-performance a-Si:H solar cells were prepared at a hydrogen-todisilane dilution ratio just below the onset of microcrystallinity. On the other hand, μc-Si:H films can be 

directly deposited at substrate temperatures (Ts) of 200-300 o

C in a PE-CVD system with a varied 

hydrogen dilution ratio, R=H2:SiH4. However, the reaction gas must be highly diluted by hydrogen such 

as 20<R<100 which leads to a very low deposition rate. Many efforts have been made to increase the 

deposition rate such as using very high frequency plasma CVD and a microwave-based CVD techniques. 

Nevertheless, a one-step high growth rate process is desirable for device applications. A deposition rate of 

~10 Å/s is standard for HW-CVD technique without using hydrogen dilution. In order to utilize the 

preparation conditions to make device-quality transition films, the correlation between the preparation 

conditions, the structure of the materials, the hydrogen configuration, and the optoelectronic properties 

need to be studied. 

II.2 Samples and Experimental Conditions 

Table II.1. Preparation conditions and parameters of the transition films prepared by HW-CVD. 

Sample ID Ts 

(°C) 

R 

H2/SiH4 

Thickness 

(μm) 

Dep. Rate 

(Å/s) 

CH 

(at.%) 

Eopt 

(eV) 

n c-Si volume 

fraction (%) 

Group I H955 250 0 0.8 4.7 12.6 1.76 3.55 0 

 H946 250 1 0.5 4.7 10.9 1.80 3.60 0 

 H948 250 2 0.7 4.7 8.9 1.78 3.45 2 

 H947 250 3 0.7 4.7 6.2 1.82 3.36 43 

 H945 250 10 1.0 3.9 4.1 1.77 3.36 82 

* T519 240 20 1.1 2.3 3.9 1.75 3.24 92 

Group II H954 150 3 0.7 4.3 9.7 1.81 - 0 

 H951 200 3 0.7 4.7 8.3 1.80 3.4 3 

 H947 250 3 0.7 4.6 6.2 1.81 3.36 41 

 H950 350 3 0.7 4.7 5.5 1.78 3.41 44 

 H952 450 3 0.7 4.4 3.8 1.74 3.53 36 

 H953 550 3 0.7 4.4 3.2 1.72 3.52 37 

Ref. H949 560 0 1.0 9.0 7.1 1.72 - 0 

* The sample was deposited in the tube system.


CH, Eopt and n represent the hydrogen content, the Tauc optical gap, and the refraction index 


Intrinsic a-S:H films were deposited by HW-CVD on c-Si substrate for IR, on rough-surface 

Corning 1737 glass for PL and Raman, on smooth surface Corning 1737 glass for transmission 

spectroscopy (TRS), constant photocurrent method (CPM) and conductivity studies. Coplanar 

Cr electrodes were evaporated on the top of the films for CPM and conductivity measurements. 

Group I films were deposited at the same substrate temperature, Ts =250 o

C, with varied R from 

7 

8 

0 to 10. Group II films were deposited at the same R=3 with varied substrate temperature, Ts, 

from 150 to 550 o

C. In addition, a reference sample was deposited at 560 oC with pure silane. 

The film preparation conditions and parameters are listed in Table II.1. Raman spectra were 

measured using an XY triple spectrograph equipped with a LN2-cooled CCD camera. The data 

were taken at room temperature using the 514.5 nm line of an argon-ion laser. The power of the 

laser was kept below the level that would thermally induce crystallization. IR spectra were 

measured at room temperature by a commercial FTIR spectrometer. The details of IR 

measurements can be found elsewhere [4]. For PL measurements, a micro-refrigerator operating 

between 80 to 300 K was used to mount the sample. A 5-mW HeNe laser with a bandpass filter 

was used for PL excitation. The penetration depth of the 632.8 nm light in a-Si:H was ~1 μm 

which is greater than the sample thickness. Between consecutive measurements, the sample was 

kept at the same position and the experimental error was < 10 meV in the experimental range of 

photon energy. The luminescence signal was dispersed by a grating monochromator and then 

collected by a liquid-nitrogen-cooled Ge detector. The lock-in technique was used with a light 

chopper at a frequency of 17 Hz. The optical absorption spectra were studied by using CPM 

combined with TRS at room temperature.

II.3 Results and Discussion 

2.3.1 Raman spectrum

The material structure was measured by Raman spectra. For device-quality a-Si:H films, the TO mode 

peaks at 480 cm-1 with a bandwidth of ~ 70 cm-1. The Raman lineshape does not change with microvoid 

density in high growth rate films. Figs. 2.1(a) and 2.1(b) show the Raman results from these two groups 

of films, respectively. The Raman result from the reference film is shown on the bottom of Fig. 2.1(a). 

Concerning the hydrogen dilution dependence, when R ≤1, the film remains in the amorphous phase 

characterized by the broad band centered at 480 cm-1. When the hydrogen dilution ratio becomes R=2, 

one can see an emergence of the band at 520 cm-1. The 520 cm-1 band grows as R continues to be 

increased. Finally, the 520 cm-1 band dominated at R=10. Fig. 2.1(b) shows the substrate temperature 

300 350 400 450 500 550 600 650

Raman Shift (cm-1)

H2:SiH4

R=10

R=0 at Ts

=560 oC

 0

 1

 2

 3

(a) Ts

= 250 oC

300 350 400 450 500 550 600 650

Raman Shift (cm-1)

Ts

= 550 oC

450

350

250

200

150

(b) R=3

Fig.2.1 Raman shifts from group I and II films with (a) varied hydrogen dilution ratio and (b) 

varied substrate temperature. The reference is shown on the bottom of Fig. 2.1(a).
9 

dependence. At Ts < 200 oC, the film remains amorphous; when Ts =200 oC, there is an emergence of the 

band at 520 cm-1. When Ts ≥250 o

C, a saturated height of the narrow band at 520 cm-1 is observed after 

the onset of microcrystallinity. Unlike the case of increasing R, there is no gradual increase of the 520 cm1 band as Ts continues to be increased. Nevertheless, using H-dilution is a necessary condition for 

crystalline transition within this temperature range. As shown on the bottom curve in Fig. 2.1(a), although 

the Ts was as high as 560 o

C, the film was still in a full amorphous phase without hydrogen dilution. 

μc-Si materials consist of the a-Si matrix, the c-Si grain and the grain boundaries (g.b.). For the best curve 

fitting we used three components: the narrow crystalline band at 520 cm-1, a Gaussian function peaked at 

480 cm-1 with a full width at half maximum (FWHM) 72 cm-1, and the Gaussian function peaked at 498-

505 cm-1 which is associated with bond dilation at grain boundaries. Notice, for these two groups of films, 

the a-Si:H band was at the same energy position with the same FWHM. It implies that the growth of the 

c-Si grains does not change the local Si-Si bonding structure of the a-Si:H matrix. The crystalline and g.b. 

volume fractions, Xc, and Xgb were estimated from Xc=(Ic+Igb)/[Ic+Igb+y(L)Ia], and Xgb= 

Igb/[Ic+Igb+y(L)Ia] where Ic, Ia, and Igb are integrated intensities of the crystalline, amorphous, and 

intermediate peaks, respectively, and y =0.9 was used. The results are listed in Table II.1. 

2.3.2. IR Absorption 

0

500

1000

1500

2000

550 600 650 700 750

R= 0

R= 1

R= 2

R= 3

R=10

Absorption coefficient (cm-1)

Wavenumber (cm)-1

(1a) Ts = 250 oC

0

50

100

150

200

250

800 900 1000 1100 1200

Wavenumber (cm)-1

(1b) Ts = 250 oC

0

100

200

300

400

500

600

700

800

1800 1900 2000 2100 2200

Wavenumber (cm)-1

(1c) Ts=250 oC

0

500

1000

1500

2000

550 600 650 700 750

Ts=150 oC

 200 

 250

 350

 450

 550

Absorption coefficient (cm-1)

Wavenumber (cm)-1

(2a) R = 3

0

50

100

150

200

250

800 900 1000 1100 1200

Wavenumber (cm)-1

(2b) R = 3

1800 1900 2000 2100 2200

0

100

200

300

400

500

600

700

800

Wavenumber (cm)-1

(2c) R=3

Fig. 2.2 The IR spectra in three frequency ranges of 550-750, 750-1200 and 1800-2200 cm-1 that cover 

the Si-H wag mode, the SiH2 bending mode, Si-O modes, and the stretch modes for (1a, 1b, 1c) group I 

and (2a, 2b, 2c) group II films. The values of R and Ts are indicated in the legend. 
cm

Hydrogen plays a crucial role in determining the structural phase transition of a-Si:H. The question then 

arises as to what are the hydrogen distribution and local bonding configuration in regard to the 

microcrystallinity. IR spectra were used to study the Si-H vibration absorption. Generally, IR spectra of aSi:H consist of three absorption regions: at 630 cm-1 related to the SiH wag mode, a doublet at 850, 890 

cm

-1 related to SiH2 bending modes, and the peaks at 2000-2090 cm-1 related to SiH stretch modes. In 

Figs. 2.2(1a, 1b, 1c) and 2 (2a, 2b, 2c) we show the IR spectra in the three frequency ranges of 550-750 

-1, 750-1200 cm-1and 1800-2200 cm-1 that cover the SiH wag mode, the SiH2 bending mode, SiO 

modes, and the stretch modes. The IR spectrum from the reference film is not shown in Fig. 2.2. Instead, 

the calculated results are listed in Table II.2. All the peaks in Fig. 2 were fitted using Gaussian functions 

and the integral intensity I=ω−1α(ω)δω was calculated. Assuming the absorption strength, A, is a constant, 

the corresponding hydrogen concentrations, N630, N850,895, N2090 and N2000, were calculated using A630 = 

-2 2 2.1x1019cm-2, A850,895 = 2x1019cm-2, A2090= 2.2x1020cm and A2000= 9.0x1019 cm . The peak energy 

positions and the bandwidth (FWHM) for the SiH wag and stretch modes as well as the percentage of the 

850-895 cm-1 mode concentration, N850,895/N630= I850,895/I630, and the percentage of the 2090 cm-1 mode 

absorption, I2090/(I2000+ I2090) are listed in Table II.2. For both group I and II films, the IR spectra mainly 

contain the monohydride modes at 630 cm-1 and at 2000-2090 cm-1. The absorption of the dihydridebending mode at 850-895 cm-1 is an order of magnitude lower compared to the 630 cm-1 absorption. 

Table II.2 The energy peak position and the bandwidth of the Si-H vibration modes. 

Sample Ts R Wag mode Stretch mode 1 Stretch mode 2 I2090/ N850,895/ 

ID Peak , FWHM Peak, FWHM Peak, FWHM (I2000+I2090) NH 

(°C) (cm-1) (cm-1) (cm-1) (%) (%) 

H955 250 0 640 89 2000 86 2085 62 32 3 

H946 250 1 640 89 1995 94 2088 62 23 3 

H948 250 2 635 87 2002 94 2088 62 38 3 

H947 250 3 631 75 2005 94 2090 62 52 5 

H945 250 10 631* 68 2010 113 2088 62 68 5 

H954 150 3 630 87 2000 112 2090 69 38 4 

H951 200 3 630 87 2000 112 2092 69 42 5 

H947 250 3 630 76 2005 95 2090 60 52 5 

H950 350 3 630 69 2001 86 2088 55 50 3 

H952 450 3 630 64 1995 86 2088 55 46 4 

H953 550 3 630 57 1996 67 2088 45 50 5 

H949 560 0 635 95 1996 86 2088 86 14 0.4 


* There is a small shoulder at 690 cm-1 

Therefore, we calculate the hydrogen concentration from the wag mode absorption. In comparison, it will 

be calculated from the stretch mode later. Assuming all the hydrogen is bonded in the matrix, the IR 

absorption at frequency ω is related to the hydrogen concentration by 

NH = AI = A∫[α(ω)/ω]dω (2.1) 

where A is sensitive to the local field. Usually, the value of A is deduced experimentally using 

independent techniques to measure the total hydrogen content. Previous work has shown that for both aSi:H and hydrogenated microcrystalline silicon, the absorption strength of the Si-H wag mode at 630 cm-1 

is independent of hydrogen content and of sample preparation. This makes using the Si-H wag intensity a 

good measure of the hydrogen concentration, which we do by 

10

N630=A630I630. (2.2) 

-3 The results of hydrogen contents, CH= N630/NSi, are listed in Table II.1. NSi= 5 x1022

 cm is used. The 

difference is <±5% between the integrated absorption peak from 550 to 750 cm-1 and the integral of the 

Gaussian function. The hydrogen content, CH, decreases monotonically by increasing either the hydrogen 

dilution ratio or the substrate temperature. We will show the detailed features of both the CH and the 

stretch modes in correlation with crystallinity later in Fig.2.3. 

0 

5 

10 

15 

0 

20 

40 

60 

80 

100 

0 2 4 6 8 

CH 

I2090/(I2000+I2090) 

c-Si volume fraction 

C 

H (at.%) 

I2090/(I2000+I2090), c-Si volume fraction (%) 

Hydrogen dilution ratio, R 

(a) 

0 

2 

4 

6 

8 

10 

12 

0 

20 

40 

60 

80 

100 200 300 400 500 600 

CH I2090/(I2000+I2090) 

Reference 

c-Si volume fraction 

C 

H (at.%) 

I2090/(I2000+I2090), c-Si volume fraction (%) 

Substrate temperature, Ts (°C) 

(b) 

H (� -10 0+I ) (�) 

y y (a) g p (b) g p

-2 10 12 

Fig. 2.3 The correlation of C ) and the relative intensity of the 2090 cm peak I209 /(I200 2090

with microcr stallinit for rou I and rou II films. 

Two previous works [5,6] show that, in the μc-Si:H deposited by PE-CVD, the frequency of the IR wag 

mode shifted to 620 or 626 cm-1; and the absorption peaks are much narrower around 626, 900 and 2101 

cm-1. They claimed that the absorption modes can be assigned to mono- and/or dihydride bonds on (100) 

and (111) surfaces in silicon crystallites. In these HW-CVD films, we have not found either the low 

frequency shift of the wag mode or a simple narrowing of the FWHM for μc-Si compared to a-Si:H. In 

Figs. 2.2(1a) and 2.2(2a) and Table II.2, one can see that the frequency of the SiH wag mode is located at 

630-640 cm-1 regardless of whether the film structure is a-Si:H or μc-Si:H. This wag-mode frequency is 

typical for standard a-Si:H. In group I films, as the crystallinity increases from 0 to ~82%, the FWHM of 

the wag mode decreases from 89 to 68 cm-1, whereas the FWHM of the 2000 cm-1 stretch mode increases 

from 86 to 113 cm-1 without a change of the FWHM in the 2090 cm-1 stretch mode. In group II films, the 

FWHM for both wag and stretch modes behaves the same, i.e. slightly decreases with an increase of 

crystallinity when 150 °C < Ts <350 °C, and rapidly decreases with an increase of the substrate 

temperature for Ts > 350 °C without an increase of crystallinity. Hence, for the R=3 films, the main factor 

that accounts for the decrease in FWHM is the substrate temperature. If one only looks at one pair of the 

films deposited at high Ts with and without hydrogen dilution, such as H949 and H953, one may find the 

FWHM of both the wag mode and stretch mode decreased for the μc-Si:H film, H953. In general, we did 

not find a simple correlation between the FWHM and the crystallinity. However, we observed a clear 

correlation between the material's crystallinity and a frequency shift of both the bending and the stretch 

modes. One can see in Fig. 2.2(1b) and 2.2(2b) that the amplitude of the 850-895 cm-1 absorption is one 

order of magnitude weaker than the SiH wag mode absorption. As shown in the last column of Table II.2, 

the relevant hydrogen concentration of N850,890 is ≤5% of the total hydrogen for both group I and II films. 

In Fig. 2.2(1b) the 850-895 cm-1 absorption seems higher for the film with R=10. Indeed, it is also about 

11

the N

~5% of NH after subtracting the contribution from SiO absorption. Notice, the last two rows in Table II.2, 

850,890/NH is as low as 0.4% in the non-hydrogen-diluted reference film H949 but 5% in the hydrogen 

diluted film H953. It raises the questions of how hydrogen dilution can bring ten times more SiH2 

incorporation into the film at such a high substrate temperature, and consequently, whether the 850-890 

cm

-1 mode in hydrogen diluted films responds to SiH2 or SiH at inner surfaces. Indeed, Itoh et al studied 

gas effusion compared to IR absorption and showed that the 850-890 cm-1 and 2100 cm-1 modes are both 

related to Si-H bonds in grain boundaries in hydrogen diluted PE-CVD μc-Si:H films. In these HW-CVD 

films, in addition to the doublet at 850-900 cm-1, the SiO related modes at 750 and ~1000-1200 cm-1 

appear when the c-Si volume fraction is above ~ 40% as shown in the curves with R = 3 and 10 in Fig. 

2.2(1b) and with Ts = 450 oC and 550.oC in Fig. 2.2(2b). The SiO absorption is largely enhanced for the 

film with a microcrystalline volume fraction of 82% (see Fig. 2.2(1b) R=10). This is because of the fast 

oxidation in air due to the porous-like structure. The porous-like structure, perhaps, is due to the high 

growth rate of c-Si grains and left the g.b. regions with much less mass density. Compared to PE-CVD aSi:H, the HW films exhibit higher structural inhomogeneity and more percentage of clustered hydrogen. 

In the two-domain model, the mass density and the structural ordering of the high dense domain in a-Si:H 

are quite close to c-Si. It is reasonable that the structure becomes porous-like if the c-Si grains grow fast 

without plasma etching out the weak bonds. Finally, the stretch mode absorption is shown in Figs. 2.2(1c) 

and 2.2(2c). First, we calculate the hydrogen concentration from the stretch mode comparing it to that 

from the SiH wag mode absorption. We found that the hydrogen concentration, N2000 + N2090, shows 

exactly the same two regimes of fast and slow decreases as shown in Figs. 2.3(a) and 2.3(b), and is about 

5% less than that calculated from the SiH wag mode, N630. This is within the accuracy of the technique. It 

implies that the SiH2 component in these films is negligible if that stretch vibration mode contributes to 

the 2090 cm-1 absorption. One exception is the film with the highest c-Si volume fraction of 82%, in 

which the sum of N2000 + N2090 is 15% larger than N630. This could be due to the decrease of the stretch 

absorption strength in the porous like structure. 

N

I

In order to find out how the hydrogen distribution changes with the crystallinity, both the hydrogen 

content and the relative integral intensity of I2090/(I2000+I2090) are compared to the c-Si volume fraction. 

The results are shown in Figs. 2.3(a) and 2.3(b) for group I and II films. The dotted lines, the solid lines 

and the open circles represent the c-Si volume fraction, the hydrogen content and the I2090/(I2000+I2090), 

respectively. At first, let us look at the solid lines that show the structural transition from a- to μc-Si as a 

function of R and Ts. One can see a rapid increase of the c-Si volume fraction at the threshold hydrogen 

dilution ratio R=2 in Fig. 2.3(a) and a threshold substrate temperature Ts = 200 ο

C in Fig. 2.3(b). The 

curves in Fig. 3(a) indicate that the c-Si volume fraction can be as large as >80% at R=10. On the other 

hand, the saturated curve in Fig. 2.3(b) indicates that, at fixed R=3, the c-Si volume fraction reaches a 

maximum of 42±1% at 250 ≤ Ts ≤ 300 οC, then slightly decreases to 33% at 350 ≤ Ts ≤ 550 οC. In a 

previous work, Kroll et al observed a considerably enhanced hydrogen concentration in/near the a/μc-Si 

transition zone in PE-CVD films. They suggested that it is due to molecular hydrogen trapped in 

microvoids. Such an enhanced CH zone has not been found in these HW-CVD films. Instead, the 

hydrogen content, CH, (see the solid dots in Fig. 2.3) decreases in two regimes of fast and slow in both 

group I and group II films. The turning points from the fast to slow regime are correlated with the onset of 

crystallinity. Based on the data from NMR and gas effusion for mixed-phase materials, most hydrogen 

clusters are located in the grain boundaries. When the material became μc-Si:H, the grain boundaries 

acted as a reservoir for hydrogen. This causes the slow decrease regime in total hydrogen content. Finally, 

the relative intensities of the 2090 cm-1 peak, I2090/(I2000+I2090), are plotted as open circles in Figs. 2.3(a) 

and 2.3(b). The big cross in Fig. 2.3(b) represents the reference film that exists in the lowest 

2090/(I2000+I2090)=14% that is equivalent to N2090/(N2000+N2090) = N2090/(N630)=30%. It is a reasonable value 

for the clustered hydrogen in an a-Si:H film deposited at high substrate temperature of 560 �C. The 

overall I2090/(I2000+I2090) changes in the range of 23-68% in group I films and 37-50% in group II films. In 

other words, the percentage of the hydrogen concentration related to the 2090 cm-1 absorption, 

2090/(N630), varies in the range of 42-83% and 60-71%. The clustered hydrogen deduced from NMR 

12

experiments is in the same range. Interestingly, the changes of the ratio I2090/(I2000+I2090) totally follow the 

manner of microcrystallinity for both group I and II films. The stretch mode doublet at 2000-2100 cm-1 

has been attributed to either a difference in frequency of a SiH oscillator in SiH and SiH2 bonding 

configurations or to a difference in the stretching frequency of SiH groups incorporated in the bulk 

material or located at internal surfaces. We examine what will be the case in these films as follows. First, 

we have examined whether there is any correlation of the 850-895 cm-1 peaks with the crystallinity. We 

found that the changes of the integrated intensity of 850-895 cm-1 peaks are not correlated to the materials 

crystallinity at all. Indeed, most of them change in an opposite way compared to the material's 

crystallinity. Therefore, the contribution from the SiH2 stretch vibration to the enhancement of the 2090 

cm

-1 absorption is negligible. 

We attribute the 2090 cm-1 absorption to the stretch vibration of clustered SiH in the g.b. regions. 

According to Eq. (2.1) we can write: 

I2090/(I2000+I2090) = N2090/[(A2090/A2000)N2000+N2090]. (2.3) 

The increase of I2090/(I2000+I2090) could be caused by a decrease of the absorption strength, A2090, and/or an 

increase of the relative concentration, N2090. We could argue as follows: a) the absorption strength A is 

sensitive to the environment of the charged oscillators so that the less dense g.b. regions could result in a 

smaller A2090. Hence, more g.b. volume fraction results in an increase of I2090/(I2000+I2090). Recent 

observation of a much smaller A2000 in a high growth rate film supports this argument. b) Straightforward, 

the relative clustered SiH bond concentration, N2090/[N2000+N2090], is higher in the g.b. regions than that in 

the a-Si:H matrix. This is supported by the NMR, gas effusion results and the slow decrease regime of the 

CH. Nevertheless, both the factors of a) and b) could contribute to the changes of I2090/(I2000+I2090). The 

next question is whether the SiH bonds are sticking on the cone-shape c-Si gain surfaces or just clustered 

in the g.b. regions. The latter is more likely because there are 32% and 14% of I2090/(I2000+I2090) in the aSi:H films without c-Si grains (see sample H955 and H949 with R=0 in Table II.2), and the changes of 

the relative intensity, I2090/(I2000+I2090), correlated with the g.b. volume fraction well (see Fig.2.3). In 

addition, the frequency range of 2085-2092 cm-1 (see Table II.2), is consistent with the SiH stretch 

frequencies of 2080-2098 cm-1 that were found in a-Si surfaces. So, we attribute the ~2090 cm-1 

absorption to clustered SiH in the g.b. regions. This is in agreement with previous work. 

2.3.3 Optical absorption spectra 

TRS and optical absorption spectra at room temperature were used to study the optical constant and the 

density of electronic states. The measurements were carried out by light incident through the top surface 

of the film. The refraction index, n, was deduced from the TRS data in the long wavelength region 

between 1500 to 2000 nm. The Tauc optical gap was deduced from the absorption spectrum between 1.7 

to 3.0 eV. The results are listed in Table II.1. We will not discuss them in detail but concentrate on the 

optical absorption spectra. The intrinsic absorption data on the high-energy side are deduced from TRS, 

whereas, CPM gives the low-energy absorption data where the penetration depth is large. Therefore, the 

absorption coefficient, α, is the average through the whole film. Figures 2.4(a) and 2.4(b) show the 

optical absorption spectra obtained by using CPM and fitted by TRS for group I and II films, respectively. 

The typical absorption spectrum of c-Si is plotted in Fig. 2.4(a) for reference. The absorption is 

contributed by the transitions from three components: the a-Si:H matrix, the c-Si grains and the grain 

boundaries. α at 1.4 eV is as small as ≈1 cm

-1 in a device-quality a-Si:H matrix. It is correlated to the 

bottom of the tail states of a-Si:H. The extra absorption coefficient at 1.4 eV shall represent the transition 

in c-Si grains and the g.b. We plot the absorption coefficient at 1.4 eV as a function of R, and Ts in Figs. 

2.5(a) and 2.5(b) to extract the information of density of states (DOS) related to microcrystallinity. 

13

14

The dotted and the solid lines indicate the c-Si volume fractions and the absorption at 1.4 eV, 

respectively. For the first group of films shown in Fig. 2.5(a), the data of α(1.4 eV) in logarithmic plot 

and c-Si volume fraction are clearly correlated. When R increases, the density of electronic states related 

to α(1.4 eV) increases due to increase of crystallinity. Meanwhile, the increase of hydrogen dilution ratio 

at Ts=250 o

C does not obviously increase the defect DOS in the a-Si matrix. This will be seen in the 

relatively large radiative recombination efficiency in PL. For the second group of films, when Ts > 250 o

C, the c-Si volume fraction does not increase but α�1.4 eV) does. This indicates an increase of both the 

tail states and the defect states in the a-Si matrix at high deposition temperature due to less hydrogen 

incorporation. This will be seen in the relatively low radiative recombination efficiency in PL, and the 

pinned Fermi energy position in the following sections.

100

101

102

103

-20

0

20

40

60

80

100

120

-2 0 2 4 6 8 10 12

 at 1.4 eV (cm-1)

c-Si volume fraction (%)

Hydrogen dilution ratio, R

α

(a)

100

101

102

103

-20

0

20

40

60

80

100

120

100 200 300 400 500 600

 at 1.4 eV (cm-1)

c-Si volume fraction (%)

Substrate temperature, Ts (oC)

α

(b)

Fig. 2.5 the optical absorption coefficient at 1.4 eV (�) and the c-Si volume fraction (∆) as a function 

of (a) R, and (b) Ts. 

10-1

100

101

102

103

104

105

0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2

c-Si

R=0

R=1

R=2

R=3

R=10

Optical absorption coefficient (cm-1)

Photon energy (eV)

(a) Ts

=250 oC

10-1

100

101

102

103

104

105

0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2

Ts = 200 oC

 250

 350

 450

 550

Optical absorption coefficient (cm-1)

Photon energy (eV)

(b) R=3

Fig. 2.4 Optical absorption spectra using CPM fitted by TRS for (a) group I and (b) group II 

films. The absorption spectra for c-Si is plotted in (a) for reference. 
2.3.4 Conductivity activation energy 

Dark-conductivity temperature dependence is also measured. The conductivity curves were obtained by 

cycling the temperature up and down. The thermally activated conductivity of σ=σoexp(-Ea/kT) was 

obtained by using the decreasing-temperature conductivity curves. We obtain the value of the 

Conductivity activation energy, Ea, by an exponential fitting. Assuming electron conduction domination, 

the value of the activation energy indicates how much the position of the Fermi level, Ef, is below the 

conduction band edge. Figs. 6(a) and 6(b) show the Ea as a function of H-dilution ratio and substrate 

temperature for group I and II films. Again, the same thresholds of the decrease of Ea and the increase of 

microcrystallinity are shown in both groups of films. The detailed correlation between Ea and the material 

microcrystallinity is clearly shown in Fig. 2.6(a). As the material became more crystallized with an 

increase of R, the Fermi energy position moved up towards the conduction band edge. This implies that 

the energy position of the states at the g.b. regions is located at the upper part of the gap. On the other 

hand, in Fig. 2.6(b) when Ts > 250 o

C, the Ea did not change much as Ts continually increased. This 

indicates that the defect states created by a high substrate temperature are located around the middle of 

the gap. Therefore, it does not change the position of Fermi level. 

0.0 

0.2 

0.4 

0.6 

0.8 

1.0 

0 

20 

40 

60 

80 

-2 0 2 4 6 8 

E a (eV)

c-Si volume fraction (%) 

R=H2:SiH4 

(a) Ts =250 oC 

0.0 

0.2 

0.4 

0.6 

0.8 

1.0 

-20 

0 

20 

40 

60 

80 

100 200 300 400 500 600 

E a (eV)

c-Si volume fraction (%) 

Ts (oC) 

(b) R=3

(� ∆ s. 

10 12 

 Fig. 2.6 Position of Fermi-level ) and the c-Si volume fraction ( ) as a function of (a) R, and (b) T

2.3.5 PL spectra at 80 K 

PL is another commonly used technique to study the DOS. In a-Si:H films, the PL signal at low 

temperatures originates from radiative band tail-to-band tail transitions. Hence, we obtain information 

about the electronic states including the optical gap, tail states and density of defects. A low energy band 

at ~ 1.0 eV has been observed in μc-Si films. For these two groups of films and the reference film, the 

relative intensity of the PL spectra at 80 K is given in Figs. 2.7(a) and 2.7(b). The bottom dotted line in 

Fig. 2.7(a) represents the PL spectrum from the reference film. The intensity of the PL is 100 times larger 

for the films with R= 0, 1, and 2 films in Fig. 2.7(a) and 50 times larger for the films with Ts = 150 o

C and 

200 o

C in Fig. 2.7(b). It means that the relative PL intensity is higher in group I films and lower in group 

II films. The PL intensity is especially low for the films with high Ts = 450 and 550 o

C. This is due to the 

high density of non-radiative centers, in agreement with the optical absorption data in Figs. 2.5(b) and 

2.6(b), whereas, more radiative centers exist at the g.b regions in the hydrogen diluted films. Comparing 

15

the top curves in Fig. 2.7(a) and 2.7(b) for the two μc-Si films, one can see the PL efficiency from the 

high hydrogen diluted μc-Si is much higher than that from the high Ts μc-Si. For the reference film shown 

on the bottom curve in Fig. 2.7(a), the relative PL efficiency is as high as the a-Si:H film deposited at 250 

o

C except the peak energy is lower. This is consistent with the Raman data shown in Fig. 2.1(a) that HW 

film deposited at high Ts without hydrogen dilution is fully amorphous. It is clear that the PL features also 

indicate a sudden change from amorphous to microcrystalline at the same thresholds of R and Ts. One 

observes that the low energy PL from grain boundaries becomes dominant and the PL total intensity 

decreases as either R or Ts increases. 

0.6 0.8 1.0 1.2 1.4 1.6 

PL relative intensity at 80 K 

Photon Energy (eV) 

(a) Ts 

=250 oC 

R=10

 0, 100 

1, 100 ×

 2, 100 × 

×

 3 

0, 100 

T s 

=560 oC 

× 

0.6 0.8 1.0 1.2 1.4 1.6 

PL relative intensity at 80 K 

Photon Energy (eV) 

× 

×

 150 o

C, 50

 200 o

C, 50

 250 o

C

 350 o

C

 450 o

C 

(b) R=3 

T s 

=550 oC 

Fig. 2.7 Relative PL spectral intensity at 80 K for (a) group I and (b) group II films. The PL 

spectrum for the reference film is show on the bottom of Fig. 1(a). 

II.4. Summary 

We studied the hydrogen incorporation and the optoelectronic properties in transition films from a- to μcSi made by hot-wire CVD as a function of hydrogen dilution of 0<R<10 and substrate temperature of 150 

°C < Ts < 550 °C. The c-Si volume fraction, Xc=(Ic+ Igb)/ [Ic+Igb+y(L)Ia], deduced from Raman 

spectra, is used to describe the degree of crystallinity. Raman results indicate that there is a threshold Hdilution ratio R=2 at Ts = 250 °C and a threshold Ts=200 °C at R=3 for the onset of the a- to μc-Si 

transition. The maximum c-Si volume fraction 82% was found with R=10 at Ts = 250 °C. We found that 

the monohydride silicon bonding is the dominant configuration in both a-Si:H and μc-Si films being 

studied. We have not found an enhanced hydrogen concentration in the a- to μc-Si:H transition zone and 

the low frequency shift of the wag mode to 626 or 620 cm-1, and there is no simple narrowing of the 

FWHM with increasing crystallinity as reported for PE-CVD films. However, we observed that both the 

hydrogen bonding distribution and optoelectronic properties are well correlated to the material's 

crystallinity. In summary: 

• With an increase in either R or Ts, the hydrogen content, CH, decreases in a fast and slow regime 

separated by the onset of microcrystallinity. We describe that this is due to the grain boundaries 

acting as a reservoir for hydrogen, which makes the decrease of total hydrogen content slow down 

with an increase of R and Ts. 

16

• The relative absorption of the 2090 cm-1 peak, I2090/(I2000+I2090), but not the 850-890 cm-1 peak, 

increases with c-Si volume fraction. Therefore, we attribute the 2090 cm-1 peak to the SiH stretch 

vibration from the clustered SiH in the grain boundary regions but not due to silicon dihydride. 

Meanwhile, the frequency of the wag mode remains unchanged at 630-640 cm-1 no matter whether 

the SiH oscillator is clustered or distributed. 

T

• The density of electronic states was studied by optical absorption spectra. The position of the Fermi 

level was measured by the conductivity temperature dependence. The radiative states were studied by 

PL spectra at 80 K. The absorption coefficient at 1.4 eV, α(1.4 eV), is used to describe the DOS 

related to microcrystallinity. We found that the same threshold of the material's crystallinity is also 

visible in the increase of α(1.4 eV), in its relevant PL(1.0 eV) band, and in the position of Ef. We 

explain that the increase of crystallinity results in an increase of the low energy electronic density of 

states including radiative and non-radiative centers. The low energy PL peak at ~1.0 eV originates 

from the radiative centers in the g.b. regions. More defects were created in the a-Si matrix when Ts > 

250 o

C. So, the Fermi level position is pinned at the middle of the gap in films deposited at R=3 and 

s > 250 o

C. 

We have shown both the features of the hydrogen incorporation and the electronic density of states are 

well correlated to the crystalline threshold and the changes of the c-Si volume fraction. We attribute the 

above features mainly related to the g.b. regions. The arguments are as follows. The hydrogen content is 

low in the tetrahedron c-Si lattice. The increase of c-Si grains would result in a decrease of the total 

hydrogen content, but does not explain the two regime of the CH decrease. Furthermore, in response to the 

2090 cm-1 absorption, the SiH bonds are not sticking on the cone-shape c-Si grain surfaces but just 

clustered in the g.b. regions as we have discussed in section III.2. In crystalline silicon, on the other hand, 

the electronic gap states are low, and the PL efficiency is limited by momentum conservation. Hence, the 

low energy luminescence centers can only be in the g.b. regions. The most defect states in μc-Si phase are 

also in the g.b regions and/or in the a-Si:H matrix for the high Ts films. Therefore, we attribute the above 

a) b) and c) observations mainly to the formation of the c-Si grain boundaries during microcrystallization. 

II.5. Conclusions: 

We conclude that making μc-Si films by the HW-CVD technique can result in a high growth rate with 

relatively low hydrogen dilution ratio. Controlling hydrogen dilution is the key factor to depositing 

device-quality μc-Si:H films. Deposition at Ts ~250 o

C with varied R produces the films with either the 

lowest defect density or the highest volume fraction of crystallinity. However, the high density of 

microvoids in the high-growth rate films cannot be detected by Raman spectra. The porous-like structure 

in the highly microcrystallized films causes the oxidation problem that can be solved. 

III. Do

 May Not Be The Dominant Recombination Centers For PC In Hot-Wire a-Si:H [7] 

III.1 Introduction 

A fundamental problem of a-Si:H is the light-induced metastable defect creation. The original StaeblerWronski paper describes that long exposure to light decreases both photoconductivity (σPC) and dark 

conductivity (σd) in a-Si:H prepared by PE-CVD. Interestingly, the σPC shows a relatively stable value 

and the σd increases upon light soaking in a-Si:H films prepared by HW-CVD. However, it was found 

that the density of metastable defects is in the same order of 1016/cm3 in both PE and HW films. As we 

have known, the typical intrinsic a-Si:H is slightly n-type and the Fermi level lies above the middle of the 

gap and the neutral silicon dangling bonds, (Do), act as the non-radiative recombination centers. 

Generally, the electron lifetime, τ, is in reverse proportion to the density of Do. The σd activation energy, 

17

Ea(A), is~ 0.7 eV at the annealed state and increased to Ea(B)~0.75 eV at the light-soaked state due to 

metastable Do creation near the middle of the gap. In order to understand the anomalous photo-induced 

changes of σPC and σd in intrinsic device-quality a-Si:H deposited by HW-CVD, we first explore the 

properties of the recombination centers by measuring thermostimulated conductivity (σTSC). 

In intrinsic a-Si:H, electron mobility is much greater than hole mobility. Using an Ohmic contact, the 

measured conductivity is attributed to electron’s motion. σTSC spectrum and temperature dependence of σPC 

have been used to explore the properties of deep states in a-Si:H. Based on the multiple-trapping model, 

Fritzsche et al found that σTSC can be written as 

σTSC = eμτ g(E)ƒ(E)1.2β⏐E⏐/T, (3.1) 

where μτ is the mobility-lifetime product, g(E) is the density of gap states, and β is the heating rate. They 

analyzed that the general shape of σTSC in a-Si:H based materials contains three regimes from ~120 K to 

~400 K: (I) an initial rise, σTSC depends on the heating rate and the Boltzmann tail of the carrier 

distribution function just before the onset of heating, ƒo(E); (II) a steady-state regime, the distribution 

function ƒ(E) is near constant and σTSC reflects μτg(E); and (III) an end of σTSC that drops to zero when all 

carriers come to thermal equilibrium. Temperature dependence of μτ(Τ) can be obtained from σPC 

temperature dependence measurements as 

σPC(T) = eμτ(Τ)G, (3.2) 

where G is the electron-hole pair generation rate. During the heating processes in both the σTSC and σPC(T) 

measurements, the change of the carrier distribution function, ƒ(E), can be described by quasi-Fermi 

level, 

Eqn(T ) = kT ln(n/Nc), (3.3) 

where n is the density of mobile carriers and Nc is the density of states at the mobility edge. On the other 

hand, the carrier distribution function just before the onset of heating, ƒo(E) can be described by a 

demarcation energy Ed that separates the thermalized and frozen-in states, 

Ed = -kT ln(νot), (3.4) 

σ

where νo ≅ 1012 s-1, and t = to is the long-delay time in σTSC experiments. Hence, the position of Ed is 

nearly fixed but the position of quasi-Fermi energy Eqn(T) changes with T during the heating process. In 

TSC measurements the carrier thermal emission rate dn/dt is balanced by the recombination rate br = 

n/τand the distribution function � ƒ(E). Generally, there are two cases of non-equilibrium carrier 

distribution: (a) For strong recombination as shown in Fig. 3.1(a), the recombination rate br is larger than 

the trapping rate, br > bt. The quasi-Fermi level Eqn is always below the demarcation energy Ed and the 

carrier distribution function ƒ(E) has a sudden jump at Ed by the reduction of mobile carrier density, n, 

due to strong recombination. (b) For weak recombination as shown in Fig. 3.1(b), br < bt. At low 

temperatures Eqn is above Ed and slowly moves down as the temperature increases finally to Eqn = Ed. The 

carrier distribution function ƒ(E) has no sudden jump in the transition zone. In the case of strong 

recombination, the theory predicts that there is a sudden rise near the transition from the initial rise to the 

steady-state regime. However, this case has never been observed in any a-Si:H-based films. Hence, the 

commonly recognized recombination centers, Do, are sort of weak recombination centers with br < bt. 

18 
log f(E) 

Electron energy E 

Eq Ed Ec 

Strong recombination, 

br > bt 

log f(E) 

Electron energy E 

Ed Eq Ec 

Weak recombination, 

br < bt 

Ed q c is the 

Fig. 3.1 Distribution function (a) for strong recombination and (b) for weak recombination relative to 

re-trapping. is the thermalization demarcation energy, E is the quasi-Fermi level and E

electron mobility edge. 

On the other hand, the commonly used techniques to measure the density of defects in a-Si:H are electron 

spin resonance (ESR) and the CPM. ESR measures the density of neutral dangling bonds Do and CPM 

takes into account all the electron transitions from gap states to conduction band edge. Whereas, a large 

discrepancy between the defect densities was found by a comparative study of ESR and the CPM in aSi:H films. For the initial state, the main cause of the difference in undoped a-Si:H was found to be the 

presence of the thin near-surface layer with a high density of DBs blind to CPM and the presence of 

negatively charged DBs inactive to ESR; for the light-soaked state, the inhomogeneous distribution of 

photocreated DBs caused by the penetration depth of the light being less than the film thickness could be 

the main contribution [8,9]. We studied the recombination centers by ESR/CPM in both HW- and PECVD a-Si:H films. 

III.2. Sample and Experimental 

Table III.1. Sample preparation conditions 

Sample 

ID 

Growth 

method 

Subs. T 

Ts (

o

C) 

Filt.current 

(A) 

Reaction gas Thickness 

(μm) 

Supplier Measurements 

T773 HW 360 16 SiH4 1.1 NREL TSC, Ea 

T779 HW 360 15 SiH4 0.45 NREL TSC, Ea 

T837 HW 400 15 SiH4 1.2 NREL TSC, Ea 

A7119 dc-PE 200 SiH4 3.0 BP-SLX TSC, Ea 

11961 rf-PE ~200 Si2H6 low H-dilut. 0.5 USSC TSC, Ea 

H954 HW 150 16 SiH4 0.7 NREL ESR/CPM 

H955 HW 250 16 SiH4 0.8 NREL ESR/CPM 

PE#1 rf-PE 250-300 SiH4 0.8 Kanazawa Univ. ESR/CPM 

19 
Device-quality intrinsic a-Si:H films were prepared by HW-CVD at the National Renewable Energy Lab 

(NREL) and by PE-CVD at BPSolarex, United Solar Ovonic Corporation (USOC) and Kanazawa 

University. The film deposition conditions are listed in Table III.1. T773, T779, T873, A7119, and 11961 

were deposited on Corning glass for σPC(T), �TSC and σd(T) studies. H954 and H955 and PE#1 were 

deposited on quartz substrates for ESR/CPM studies. The HW films of H954 and H955 were using a 

separated deposition system and deposited at a relatively low temperature compared to T773, T779 and 

T873. Their σPC were not as stable as those deposited at high temperatures, but they still were more stable 

than PECVD films and their initial activation energy was large. In σTSC and conductivity studies, the MMR 

micro-probe system with vacuum 10-4 Torr was used to mount the sample. The σPC was generated by a 

632.8 nm laser at a rate of G = 5×1015 photons/cm3

s

1

 and its temperature dependence was measured from 

80 K to 400 K. In the σTSC study the sample was first cooled down to 120 K and then exposed to the 632.8 

nm laser light with an exposure time te = 60 sec. After a delay time, to = 30 min σTSC was measured in a 

step-heating process with a constant heating rate β=1.5 K/min in the dark. Ohmic contact and a low field 

of 103 V/cm were used to avoid electrode effects. A Keithley 6512 programmable electrometer was used 

to record the current. A LabView program was used for experimental control and data acquisition. Light 

soaking was performed in situ. The sample temperature was kept at 35 oC during light soaking. State A 

was reached after the sample was annealed in a vacuum for one hour at 160 oC; state B was reached by a 

200 mW/cm2 white light illumination through a water filter of 4 hours. The details of ESR/CPM 

measurements can be found elsewhere [8,9]. 

III.3. Results 

10-12 

10-10 

10-8 

10-6 

10-4 

0 5 10 15 20 25 

PE-CVD 

HW-CVD

Conductivity (Ohm-cm)-1 

Light soaking time (h) 

A 

A 

B 

B 

Fig. 3.2 The conductivities before, during and after light 

soaking for typical intrinsic a-Si:H films made by HWCVD (solid line) and PE-CVD (dotted line). 

Figure 3.2 shows the typical photodegradation of σPC and σd for device quality 

intrinsic PE-CVD a-Si:H film deposited at 

200 oC and HW-CVD film deposited at 320 oC < TS < 440 oC. The dotted line represents 

the light-soaking effect for PE-CVD film 

that shows a decrease of both σPC and σd. 

Compared to the PE-CVD film, one finds 

the better stability of σPC and the increase of 

σd upon light soaking for the HW film. To 

explain the anomalous situation for the HWCVD films, we explore the nature of the 

defects as follows. 

3.3.1 Natural degradation effect 

The Fermi-level position is deduced from σd = σooexp(-Ea/kT). We mounted the film on a thermal stage 

in a vacuum and kept it at 120 oC for 30 min to cleanup the film surface adsorption, then we took the 

conductivity data from a temperature cycle of 270 K to 420 K to deduce the activation energy at the initial 

state. However, the value of Ea was driven down with the storage time or a temperature cycle without 

light soaking. For example, we observed Ea(I) ~0.98 eV and Ea(I) ~ 0.89 eV after being stored in the dark 

20

for several weeks for the film T773; and Ea(I) ~ 1.05 eV and Ea(A) ~ 0.98 eV after annealing at 160 oC 

for 1 h for sample T779. We named this "natural degradation" effect because of the similarity to the 

photo-degradation effect in conductivities. The values of Ea(I), and of Ea(A) and Ea(B) measured before 

and during the time period of σTSC studies are listed in Table II. One finds that the activation energy Ea 

increased from 0.73 to 0.77 eV for the non-hydrogen diluted PE-CVD film #7119. Whereas, all the HWCVD and the hydrogen-diluted PE-CVD samples show the opposite direction, i.e. the Ea decreased after 

light soaking. 

Table III.2. Conductivity activation energies at initial- annealed- and light soaked states. 

Sample ID 11961 A7119 H954 H955 T773 T779 T837 

Ea(I) (eV) - - 0.86 0.95 0.98 1.05 0.96 

Ea(A) (eV) 0.92 0.73 0.84 0.93 0.89 0.93 0.86 

Ea(B) (eV) 0.85 0.77 0.81 0.85 0.84 0.86 0.82 

3.3.2 Strong recombination center in a-Si:H prepared by HW-CVD 

10-11 

10-10 

10-9 

10-8 

10-7 

10-6 

State A 

State B 

Photoonductivity (Ohm-cm)-1 

(a) 

Dark conductivity 

10-11 

10-10 

10-9 

10-8 

10-7 

10-6 

State A 

State B 

Photoonductivity (Ohm-cm)-1 

(b) 

Dark conductivity 

2 4 6 8 10 12 2 4 6 8 10 12 

1000/T (1/K) 1000/T (1/K) 

Fig. 3.3 Temperature dependence of photo- and dark conductivity before and after 200 mW/cm2 

white light soaking for 4 h, (a) for HW-CVD film T779 and (b) for PE-CVD film A7119. 

Fig. 3.3 shows both the σd(T) and the σPC(T) curves for typical device-quality (a) HW sample T779 and 

(b) PE-CVD a-Si:H sample A7119. The thermal quench valley in all the σd(T) curves, perhaps, resulted 

from the trapped holes in the valence band tail. The additional convex peaks at ~200 K and ~255 K in 

Fig. 3.3(a) are correlated to the σTSC peaks in Fig. 3.4(a). Figs. 3.4(a) and 3.4(b) show the σTSC curves for 

thesame samples of T779 and A7119 at States A and B. We are not going to analyze the σTSC spectra in 

detail but only pay attention to the transition near 200 K. Comparing Fig. 3.4(a) with 3.4(b) the most 

interesting feature is that a sudden rise occurs in the σTSC near the transition from the initial rise to the 

steady-state regime in the HW film at both States A and B as shown in Fig. 3.4(a). According to 

Fritzsche's theory, this is an evidence of strong recombination centers in the HW a-Si:H film at both 

States A and B. It means that the recombination rate for electrons, br is larger than the trapping rate, br > 

bt. D+ that has a larger capture rate for electrons than Do can be the candidate of the strong electron 

recombination center. A sudden rise in the σTSC near 200 K has also been observed in the other HW films 

of T837 and T773; whereas, a smooth transition of σTSC similar to that in A7119 was observed in PE-CVD 

21

film 11961. Hence the sudden rise feature is the characteristic for the HW-CVD a-Si:H deposited at high 

substrate temperatures. 

10-14 

10-13 

10-12 

10-11 

10-10 

10-9 

10-8 

10-7 

State A 

State B 

TSC (Ohm-cm)-1 σ 

(a) 

10-14 

10-13 

10-12 

10-11 

10-10 

10-9 

10-8 

10-7 

State A 

State B 

TSC (Ohm-cm)-1 σ 

(b) 

2 3 4 5 6 7 8 2 3 4 5 6 7 8 

1000/T (1/K) 1000/T (1/K) 

Fig. 3.4 Thermostimulated conductivity for the same samples in Fig. 3 at annealed State A and lightsoaked State B, (a) for HW-CVD film T779 and (b) for PE-CVD film A7119. 

3.3.3 ESR and CPM studies 

1016 

1017 

Nd(A) 

Nd

Density of defects (cm

(B)

-3) 

Sample 

(a) 

H954 H955 PE#1 

Do(A)bulk 

Do(B)bulk 

ID 

(b) 

H954 H955 PE#1 

1016 

1017

 ESR, ∆Do 

CPM, ∆Nd

Density of metastable defects, 

∆Nd, (cm -3) 

Sample ID 

(c) 

H954 H955 PECVD#1 

2 Fig. 3.5 Total density of defects for HW- and PE-CVD films before and after 200 mW/cm white 

light soaking for 30 h. measured by (a) CPM, (b) ESR, and (c) light-induced metastable defects 

measured by CPM and ESR, respectively. 

We have measured the density of defect states by ESR and CPM on two HW-CVD films deposited at 

substrate temperatures of 150 and 250 oC in comparison to the PE-CVD film. Figs. 3.5(a) and 3.5(b) 

22 
show the CPM and ESR results at State A and State B (that was reached by light soaking using a 200 

mW/cm2 white light for 30 h). For the ESR data, only the bulk defect density was counted by assuming 

that the densities of the surface defect are the same for all the films based on the previous experimental 

results. Regardless of the absolute values, the light soaking results in an increase of defect density from 3 

to 9x10 16/cm 3 for all the films. The CPM data show a parallel increase of Nd, but the ESR data show a 

much larger increase of Do in the PECVD film than in the HW films. We plot the density of light-induced 

metastable defects, ∆Nd= Nd(B) - Nd((A) and ∆Do = Do(B) - Do(A), in Fig. 5(c). The interesting thing is 

that in the HW films the ∆Nd deduced from CPM is larger than ∆Do; but in the PECVD film the ∆Nd is 

smaller than ∆Do. The latter is commonly observed in PE-CVD film. Reasonably assuming the same 

inhomogeneous distribution of photocreated DBs in all the films after the same light soaking processes, 

the different results between HW- and PE-CVD films can be explained by the opposite direction 

movement of the Fermi-level upon light soaking; i.e. after light soaking the Fermi-level moved upwards 

in the HW films which left more negative charged defects, D-

, that are inactive to ESR. In the PE-CVD 

film, after light soaking the Femi-level moved downwards, which results in more Do. This means less 

metastable electron recombination centers, Do, hence less light induced σPC degradation in HW films than 

that in PECVD film. But, there are more metastable defects in their negative charged state in HW films. 

III. 4. Summary and Discussion 

3. 4.1 Summary 

In HW-CVD a-Si:H films deposited at substrate temperature at 320 oC < TS < 440 oC, we have observed a 

relatively stable photoconductivity and upward movement of the Fermi-level after light-soaking. The 

unusual photo-degradation curve for HW-CVD films in Fig. 3.3 can be explained by either a) or b) as 

follows. a) The upward movement of the Fermi-level position: Assume Do is the dominant recombination 

center and the electron lifetime is in a reverse proportion to the density of Do. The creation of metastable 

DBs could be balanced by the upward movement of the Fermi-level, which results in filling up partial Do 

into D-

. Therefore, the total number of Do and then the electron lifetime could change much less than the 

case of the Fermi-level moving downward. The CPM/ESR results in Fig. 5 support these arguments. b) 

The existence of stronger electron recombination centers such as positive charged defects, D+. We have 

shown evidence of strong recombination centers for electrons by thermal stimulated current 

measurements. If the D+ rather than Do act as dominant non-radiative recombination centers at both States 

A and B, the σPC should be relatively stable upon light-induced Do creation. As one knows, the capture 

rate of D+ for electrons is much greater than that of Do. A density of 1015/cm3

 of the D+ among total 

defects of 1016/cm3 would dominate the electron lifetime. Obviously, an increase of metastable Do on the 

order of 1016/cm3 cannot change the electron lifetime much. One cannot exclude that both a) and b) 

could play a role in the unusual photo-degradation curve for HW-CVD films. 

The comparative ESR/CPM studies show the light-induced metastable defects in the order of 1016/cm3 for 

both the HW and PE-CVD a-Si:H. The density of defects in the HW films is higher than the PE-CVD 

film in State A, but the change of the density of neutral dangling bonds, ∆Do, with light soaking is smaller 

in the HW films than the PECVD film. Whereas, the ∆Nd deduced from sub gap absorption is larger than 

∆Do in the HW films; the ∆Nd is commonly smaller than ∆Do in PECVD film. We explained this by the 

opposite direction movement of the Fermi-level in HW- and PE-CVD films upon light soaking. It is a 

blemish of this work that we do not have films T773, T779 and T837 deposited on ESR quartz substrates 

to measure ESR. 

3. 4.2 Natural degradation effect 

23 
We have observed a natural degradation effect, i.e. the activation energy is driven down with storage time 

and/or temperature cycle without light soaking. In other words, the Fermi level position moved upward 

with prolonged storage time. This movement results in an increase of the dark conductivity, an increase of 

the electron mobility-lifetime product, μnτn, and a decrease of the hole mobility-lifetime product, μπτπ, in 

other words, a decrease of the charge collection length that is crucial for a-Si:H solar cell performance. 

Therefore, a natural degradation of the a-Si:H solar cells made by HW-CVD can be understood. One may 

wonder if this degradation was caused by oxidation. This can be distinguished experimentally by storing 

the sample either in high vacuum or oxygen environment. Results indicate that it was not the oxidation 

effect, at least, for the films studied in this work. 

3. 4.3 Why the activation energy is very large and driven down? 

How can the activation energy, Ea , be as large as ~1.0 eV in a ~1.7 eV optical gap material? Why is it 

driven down with storage time, temperature cycle and light soaking? We try to understand these by an 

over simplified picture. If there is only the silicon-dangling bond as the defect, one possible explanation 

for the large Ea is a large deformation potential in the HW-CVD films. The thermal excitation energy of 

an electron consists of two terms: 

Ea = ET +W = ET +B2

/4A, (3.5) 

where ET is the energy of the defect state below the bottom of the conduction band, Ec, W is the lattice 

relaxation energy at the defect, B is the deformation potential representing the strength of the electronphonon interaction, and A is the strength of the network vibrations. If W could be as large as ~0.3 eV, one 

could obtain a value of Ea ~1.0 eV for a defect located 0.7 eV below Ec. However, the assumed lattice 

relaxation energy may be too large. The lattice relaxation energy at the defect, W, was found not to be 

larger than 0.1 eV in n-type PE-CVD a-Si:H. It is possible that U has a distribution of values so that a 

small fraction of defects may have negative U, 

U= e2/4πεεor - W < 0, (3.6) 

where e2/4πεεor is the Coulomb interaction energy. A higher probability of negative U defects exists in 

HW film since there is larger lattice relaxation energy than that in PE-CVD film. Whereas, the majority of 

defects must be positive U defects with ESR signal and unpinned Fermi level. The negative U defects 

result in the same number of D-

 and D+. That small number of D+ could be the strong recombination 

centers for electrons and the D-

 located in the deep of the gap could result in a large Ea. Alternatively, if 

modeling the disorder as potential fluctuations and the potential fluctuations exceed U; this permits 

coexistence of Do, D-

 and D+ without negative U, whereas, the large value of the initial Ea can't be 

explained. 

Besides, the strong electron-phonon interaction in HW films shall result in a larger energy difference 

between the absorption and emission in optical transitions. This was evidently shown in a lower PL peak 

energy in HW films than that in PE-CVD films; consequently, the value of Voc is lower. Finally, the 

external factors of storage time, temperature cycle and light soaking may result in a decrease of W, 

consequently Ea is driven down. Finally, the defect-complex caused by the contamination by the hot 

filament is not excluded as a candidate of the strong recombination center for electrons. It needs to be 

studied. 

III.5 Conclusions 

24 
The features of light soaking effects in a-Si:H prepared by HW-CVD are different from those in PE-CVD 

films: the PC decreases only slightly and the dark conductivity increases. We explore the properties of the 

recombination centers using ESR, CPM and σTSC. We observed a sudden rise of σTSC near the transition 

from the initial rise to the steady-state regime of σTSC. This is an evidence of strong-recombination. When 

the strong recombination centers such as D+ rather than Do dominate the electron lifetime, the PC would 

show a relatively stable value. 

IV. PL study of growth rate dependence on solar cells made by BPSolar[10,11] 

IV.1. A special microstructure of high growth rate films 

480 cm-1 

a-Si:H intrinsic films and their solar cells were 

prepared using DC-PECVD with growth rates of 1, 

3, and 10 Å/s by BPSolar. Fig. 4.1 (Raman shift) 

shows an identical TO mode for all the samples, 

which indicates the local silicon-bonding 

configuration does not change with the growth rate 

from 1 to 10 Å/s. IR results show that the hydrogen 

bonding configuration is monohydride (Si-H) 

dominated, and the hydrogen content obtained from 

Intensity (a.u.) 

the 630 cm-1 wagging mode is 12-14 at.%. 

Surprisingly, a very small absorption strength for the 

stretching 2000-2100 cm-1 mode was found for the 

samples deposited at the higher growth rates (3, 10 

Å/s). For these same samples, the PL spectra exhibit 

a red shift. Both the IR and PL results might be 250 

L0276

L0277

L0278 

70 cm-1

300 350 400 450 500 550 600 650 

related with the same microstructures formed at high Raman shift (cm-1) growth rates. However, the cell performances were 

comparable. The i-layer thickness, optical gaps and 

the solar cell performance are listed in Table IV.1. Fig. 4.1 Raman shift for samples with growth 

rates of 1, 3 and 10 Å/s. 

Table IV.1. Sample preparation conditions, i-layer properties and their solar cell performances. 

Sample Growth rate 

(Å/s) 

Thickness 

(Å) 

Optical gap 

(eV) 

FF Voc 

(V) 

Jsc 

(mA/cm) 

Eff. 

(%) 

*L171-2C 1 ~2500 0.731 0.895 11.64 7.62 

*L172-5C 3 ~2500 0.721 0.908 11.68 7.65 

L1221-2TR 10 ~2500 0.717 0.890 11.51 7.34 

*L0172-3 1 ~6000 

*L0172-4 3 ~6000 

*L0173-1 5 ~6000 

L0276-1, 1c 1 ~6000 1.711 

L0277-1, 1c 3 ~6000 1.721 

L0278-1, 1c 10 ~6000 1.715 

* These samples were deposited using a discontinued deposition-chamber configuration 

4.1.1 A special case of the IR spectra: 

25 
The IR spectra are shown in Fig. 4.2 for the samples (a) L0172-3, L0172-4 and L0173-1, and (b) L0276-

1c, L0277-1c and L0278-1c. The absorption at ~630 cm-1 was used to calculate CH using A630=2.1×1019 

cm

-2. CH=12-14 at.% was found and listed in Table IV.2. Interestingly however, the stretch modes were Absorption coefficient (cm-1) 

2000


1500


1000


500 

0 

400 800 1200 1600 2000 

Wavenumber (cm-1) 

L0172-3 

L0172-4 

L0173-1 

(a) 

L0276-1c 

L0277-1c 

L0278-1c 

0 

500 

1000 

1500 

2000 

2500 

3000 

3500 

Absorption coefficient (cm -1 )

(b) 

400 800 1200 1600 2000 2400 

Wavenumber (cm-1) 

Fig. 4.2 IR spectra for samples of (a) L0172-3, L0172-4, L0173-1, (b) L0276-1c, L0277-1c, and L0278-1c. 

unusual in the second group of films as shown in Fig. 4.2(b). The larger 2000-2100 cm-1 absorption band 

indicates the absorption strength for the stretching mode becomes much smaller than previous results. 

One can see that the stretching absorption increases incrementally when the growth rate increases from 1 

to 10 Å/s. The values of (I2000+I2090)/I630, α=(I2000/I630), and β= I2000/I2090 are listed in Table IV.2. The 

results are from the second measurements after one year. One can see that the ratio of (I2000+I2090)/I630, 

somewhat is normal for the first group of films L0172-3, L0172-4, and L0173-1 with growth rates of 1, 3 

and 5 Å/s. Whereas, the value of α for the second group of films is much larger. Especially, for the films 

with growth rate ≥ 3 Å/s the α values are about a factor of 3-4 larger than previous results. The values of 

β are also larger than previous results. One possible explanation for the large α and β is that a large 

26

27

fraction of the Si-H bonds are located at/near some microstructures as clustered monohydrides in the 

films. The samples were retained in the lab at room temperature for a year and the IR spectra were remeasured twice. The unusually large stretch modes were confirmed. 

Table IV.2. Hydrogen content and the absorption strengths of Si-H vibration modes

Sample 

ID 

CH

(at.%) 

NH

(1021

cm-3) 

I630 

(cm-1) 

I2000 

(cm-1) 

I2090 

(cm-1) 

(I2000+I2090)/I630 α=I2000/I630 

=A630/A2000

β=I2000/I2090 

=A2090/A2000

*L0172-3 12.7 6.4 302 36.2 5.1 0.15 0.13 7.1 

*L0172-4 14 7.0 333 31.2 10.3 0.13 0.10 3.0 

*L0173-1 14 7.0 333 37.2 11.2 0.16 0.12 3.3 

L0276-1c 12.3 6.4 292 76.6 13.1 0.31 0.27 5.9 

(redo) 12.4 6.2 295 90 18.7 0.39 0.32 4.8 

L0277-1c 12.2 6.1 290 148 44.8 0.65 0.50 3.3 

(redo) 13.1 6.6 311 143 41.1 0.61 0.47 3.5 

L0278-1c 13.4 6.7 319 157 44.8 0.69 0.54 3.5 

(redo) 11.7 5.9 278 179 56.0 0.87 0.66 3.2 

4.1.2 A special case of the PL spectroscopy: PL spectrum and its temperature dependence were studied 

for these two groups of samples of L0276-1, L0277-1, L0278-1, and L0171-2c, L0172-5c and L1221-

2TR. We found that unlike the IR results, both the spectral lineshape and the total intensity temperature 

dependence showed similarity between these two groups of samples. Hence, we show the results for one 

group of films. Figs. 4.3(a), 4.3(b) and 4.3(c) show the PL spectra at temperatures 80 K, 160 K and 180 

K; and Fig. 4.3(d) shows the PL peak energy, as a function of the measurement temperature, for samples 

L0276-1, L0277-1, and L0278-1. One can see that there are differences in the spectral lineshapes between 

the low growth rate film L0276-1 and the two higher growth rate films L0277-1 and L0278-1. The solid 

lines in Figs. 4.3 show that the low growth rate film L0276-1 exhibits PL characteristics that are typical 

for device quality intrinsic a-Si:H, i.e., the peak energy at ~1.4 eV at 80 K and it gradually decreases as 

the temperature increases due to carrier thermalization in tail states. However, as shown by the cross and 

dotted symbols for the growth rate = 3 Å/s films, the PL peak energy is red-shifted. There is a PL band at 

~1.1 eV in addition to the 1.4 eV band as shown in Figs 4.3(b). 

The PL peak energy position as a function of temperature is shown in Fig. 4.4. The PL peak position 

decreases rapidly at ~ 170 K due to a weak-temperature dependent PL band at ~1.1 eV. A nanovoid0.8 0.9 1.0 1.1 1.2 1.3 1.4 1.5 1.6

L0276-1

L0277-1

L0278-1

Normalized PL intensity (a.u.)

Photon Energy (eV)

(a)

0.8 0.9 1.0 1.1 1.2 1.3 1.4 1.5 1.6

Normalized PL intensity (a.u.)

Photon Energy (eV)

(b)

Fig. 4.3 PL results for samples L0276-1, L0277-1, and L0278-1 (a) spectra at 80 K, (b) spectra at 180 K. 
related large red shift of PL peak energy in hot-wire a-Si:H has been observed recently. It is possible that 

the 1.1 eV PL in the high-growth rate films is related to the same microstructures that causes a lower 

absorption strength for the Si-H stretching mode as shown in Fig. 4.2. 

1.05 

1.10 

1.15 

1.20 

1.25 

1.30 

1.35 

1.40 

L0276 

L0277 

L0278 

PLpeak (eV) 

(d) 

50.0 100.0 150.0 200.0 250.0 300.0 

Temperature (K) 

Fig. 4.4 PL peak energy as a function of temperature 

for samples L0276-1, L0277-1, and L0278-1. 

4.1.3. Summary of the special case: The IR and PL results indicate that certain microstructures exist when 

the growth rate is ≥3 Å/s. In the following section we report the results from a second group of samples 

made at BPSolar using a modified deposition system. The above IR and PL results shown in Figs 4.2 and 

4.3 were not reproducible in the second group of samples although the same growth conditions were used. 

We did not find the microstructure related unusual results of IR and PL. It indicates that the sample 

quality depends on not only the growth conditions but also the deposition system. 

IV.2. PL Results from a Group of New Samples 

4.2.1. Sample and Experimental: Four pairs of BP-Solar films, L2127, L2129, L2136, and L2137 at 

original State A and light-soaked state B (exposed to 1 sun, Na vapor lamp for 600 h at 50 o

C) were 

studied by PL spectroscopy. We have also done the measurements after in-situ light soaking using a 200 

mW/cm2

 white light at 40 oC for two hours. The sample conditions and their PL parameters are listed in 

table IV.3. The PL energy spectra were measured at the temperature range of 80 K - 300 K 

Table IV.3 The sample preparing conditions and their PL parameters 

Sample ID Growth rate 

(Å/s) 

Thickness 

(μm) 

Peak main at 80 K

 (eV) 

Slope of Imain, TL 

(K) 

L2127 1 1.79 1.35 24 

L2129 3 1.47 1.36 27 

L2136 5 1.65 1.36 28 

L2137 10 1.76 1.34 40 

where Peakmain and Imain are the peak energy position and the integral intensity of the PL main band. 

28 
29

4.2.2. PL Results 

PL energy spectra at low temperatures. Figure 4.5 shows the PL spectra at 80 K for those four samples 

listed in table IV.3. The spectral lineshapes show the same features as that of typical device-quality 

intrinsic a-Si:H; whereas, the peak energy position, Peakmain, is slightly different. The highest growth rate 

film shows the lowest Peakmain that can be related to a broader band tail and/or a narrower optical gap. 

Fig. 4.6 shows the PL spectra at 80-160 K for sample L2127. One can see that both the PL intensity and 

the Peakmain, decrease when the temperature increases due to carrier-themolization in the band tail states.

0.8 0.9 1 1.1 1.2 1.3 1.4 1.5 1.6

L2127/80 K 

L2129/80 K

L2316/80 K

L2317/80 K

Normalized PL intensity

Photon Energy (eV)

Fig. 4.5 PL spectra of samples with growth rate 

of 1, 3, 5 and 10 Å/s at 80 K.

0

200

400

600

800

1000

1200

1400

0.6 0.8 1 1.2 1.4 1.6

80 K

90 K

100 K

120 K

140 K

160 K

Relative PL intensity (a. u.)

Photon Energy (eV)

Fig. 4.6 PL spectra of sample with growth rate 

of 1 Å/s at temperatures of 80 - 160 K. 

1.1

1.2

1.3

1.4

50 100 150 200 250 300

L2127

L2129

L2136

L2137

PL peak energy of the main band, Peakmain (eV )

Temperature (K) 

10-2

10-1

100

101

102

50 100 150 200 250 300

LI2127, TL ~ 24 K

L2129, 27 K

L2136, 28 K

L2137, 40 K

Relative PL intensity of the main band, Imain (eV )

Temperature (K)

Fig. 4.7 Peakmain of samples L2127, L2129, L2316 

and L2317 as a function of temperature.

Fig. 4.8 Imain of samples L2127, L2129, L2316 

 and L2317 as a function of temperature.
I

PL intensity temperature dependence: The Peakmain position and the integral intensity, Imain, as a function 

of temperature at State A for those four films are plotted in Figs. 4.7 and 4.8, respectively. One can see 

that for the samples with a growth rate from 1 to 5 Å/ both Peakmain(T) and Imain(T) is similar as that for 

typical device-quality a-Si:H. However, for the sample with a growth rate of 10 Å/s both the Peakmain and 

main show much weaker temperature dependence. Especially, the PL peak energy is "saturated" at ~ 1.25 

eV when T > 160 K. This indicates a deep trapping state in this high growth rate sample. The slope of the 

Imain, TL, is listed in table IV.3. 

4.2.3 In-situ light soaking: The Imain decreased slightly and the integral defect PL intensity, Idefect, clearly 

increased after light soaking for two hours. Figure 4.9 shows the Idefect as a function of T at State A for 

those four samples. Figure 4.10 shows the light-induced increase of Idefect at 260 K. 

100 

101 

160 180 200 220 240 260 280 300 320 

L2127 

L2129 

L2136 

L2137 

Relative PL intensity of the defect band, Idefect (a. u.) 

Temperature (K) 

0 2 4 6 8 

State A 

State B 

0.4 

0.6 

0.8 

1.0 

1.2 

1.4 

1.6 

Idefect at 260 K (a.u) 

Growth rate (Å/s) 

I

p p g g g. 

10 12 

Fig. 4.9 Defect PL intensity Idefect as a function Fig. 4.10 defect at 260 K as a function of the 

of tem erature for the four sam les. rowth rate before and after li ht soakin Relative PL intensity (a.u.) 

104 

103 

102 

101 

100 

Imain(A) 

Imain(B) 

Idefect(A) 

Idefect(B) 

Peakmain(A) 

Peakmain(B) 

the light soaking was not in-situ, the tendencies of 0.8 

light-induced changes are the same, i.e. a decrease 

10-1 0.7 of the Imain and an increase of the Idefect. 50 100 150 200 250 300 350


Temperature (K)


30 

1.4 

Fig. 4.11 the Peakmain, Imain, and Idefect vs. T for 1.3 

sample L2127 at State A and State B.


1.2 

1.1 

1.0 

PL peak energy (eV) 

Finally, we show the Peakmain, Imain, and Idefect for 

sample L2127 at State A and State B (50 o

C, 1 sun, 0.9 Na vapor lamp for 600 h) in Fig. 4.11. Although 
IV.3. Summary and Discussion 

Unlike the unusual IR and PL results of the first group of samples, the results from the second group of 

samples are usual (because the deposition system was modified). It suggests that increasing the deposition 

rate from 1 to 5 Å/s has a modest effect on the electronic tail states and a gradual increase of defect states. 

The latter, perhaps, will not affect the solar cell performance. Especially, the features of the samples with 

a growth rate of 3 and 5 Å/s are very close to each other. An increase of the growth rate to 10 Å/s has a 

dramatic effect in the electronic states. In additional to the broader tail states and higher density of 

dangling bond states, the deep-trapping states may be related to the microvoids. The material is much 

more defective compared to the others, that is in consisting with the results of the microstructural studies 

[11]. The total PL intensity decreased and the relative defect PL intensity increased after light-induced 

metastable defect creation. We have noticed that the low-growth-rate film (1 Å/s), L2127, shows some 

unusual features such as the peak energy position is lower than that of the films with a higher growth rate 

of 3 and 5 Å/s (se Figs. 4.5 and 4.7) and an increase instead of a decrease of both the Peakmain and Imain 

after light soaking (see Fig.4.11). We have double-checked the data that were reproducible. The reason 

remains unclear. 

V. Systematical Raman study of thin films from amorphous-to-nanocrystalline silicon prepared by 

HW-CVD [12] 

V.1 Introduction 

Raman spectroscopy is a sensitive tool to study the structures of a-Si-based materials. In a-Si, all phonon 

modes of the transverse acoustic (TA), longitudinal acoustic (LA), longitudinal optical (LO), and TO 

modes are Raman active. Thus, Raman spectra from a-Si at room temperature yield a reasonable spectral 

comparison with phonon density of states that are modified substantially by small changes in short-range 

order (SRO). Theoretical calculations found that the width FWHM, the peak frequency of the TO phonon 

mode (ωTO), and the relative intensity of the TA mode, TA/TO, are sensitive to bond-angle variation (∆Θ) 

in an a-Si network. Beeman et al. [13] found a linear relationship between the width of the TO mode and 

the spread in the mean bond angle ∆Θ in an a-Si network as 

FWHM = 15+6∆Θ. (5.1) 

More recent calculations using 1000-atom configurations showed a similar relation of Γ=18.4 + 6.6∆Θ 

and a shift of the TO mode frequency toward higher frequency as ∆Θ decreases following 

ωTO = -2.5 ∆Θ +505.5 (5.2) 

for HV-polarized light. The authors further found that the TA/TO intensity ratio decreases linearly with 

decreasing ∆Θ. For device-quality a-Si:H, the TO mode centered at 480 cm-1 with FWHM ~70 cm-1. 

According to equation (5.1), the spread in mean bond angle ∆Θ is ≤10°. In c-Si, on the other hand, the 

Raman spectrum consists of a single sharp TO mode with FWHM of 4 cm-1 at a frequency of 520 cm-1. 

All other modes are not Raman active because of the symmetry in the fcc lattice. However, when the c-Si 

grain size is as small as a few nm, the momentum conservation will be relaxed and Raman active modes 

will not be limited to the center of the Brillouin zone. Thus, the frequency could shift from 520 cm-1 to 

~512 cm-1 with decreasing grain size from ≥10 to ~3 nm. Below 3 nm, a crystalline-to-amorphous 

transition occurs. In addition to the a-Si and c-Si phonon modes, the Raman TO mode from nc-Si films 

contains a third intermediate peak centered around 500 ± 10 cm-1 which could be attributed to the 

contribution of grain boundaries or other structures. With hydrogen dilution, Tsu et al. found chainlike 

31

objects in PE-CVD films using high-resolution transmission electronic microscopy (TEM) which they 

claim is evidence for intermediate ordering in H-diluted films. They found that the a-Si TO band centered 

at 490 cm-1 (37 ± 3 cm-1 full width) instead of 480 cm-1. 

To determine the c-Si volume fraction from the Raman TO mode, one must take into account the 

differences between a-Si and c-Si with respect to scattering cross sections and optical absorption 

coefficients. Bustarret et al. found that the ratio of the cross section for the amorphous-to-crystalline 

phase for grain size > 3 nm at the excitation wavelength 514.5 nm can be written as

 y(L) = 0.1 + exp[-(L/250)]. (5.3) 

It is expected that y(L) will be 2.4 times smaller at 632.8-nm excitation. Based on the above argument, we 

deduced the crystalline volume fraction, Xc, from the Raman TO mode using 

Xc =(Ic + Igb)/[Ic + Igb +y(L)Ia], (5.4) 

where Ic + Igb and Ia are integrated intensities of the c-Si, intermediate, and a-Si peaks, respectively; y(L) 

decreases from 1 to 0.4 when the grain size L grows from 3 nm to 30 nm. We found that when using 

514.5-nm laser front excitation, Xc can be deduced in such a way to be consistent with the results from xray diffraction (XRD) for HW-CVD transition films. In previous works, we reported Raman, XDR, FTIR, 

PL, and optical absorption results for HW-CVD films with either varied hydrogen-to-silane ratio, R, with 

fixed silane flow rate or varied substrate temperature, Ts, at fixed R = 3. We observed a blue shift of the 

a-Si:H PL peak frequency, which was due to narrowing of the band tail states by hydrogen dilution and a 

low-energy PL band in the nc-Si films. The low-energy PL band originates from the c-Si g.b. regions, and 

thus, we have attributed the intermediate component of the Raman TO mode to the g.b. regions. The 

previous Raman results showed a threshold of R ≈ 2 at Ts = 240ο

-250οC and Ts ≈ 200οC at R = 3 for the 

structural transition from a- to nc-Si:H. Both the threshold growth conditions of crystallinity and the 

crystalline silicon volume fraction deduced from the Raman TO modes are consistent with the results 

from XRD. In this work, we have extended our deposition parameter range to include not only the Hdilution ratio R, but also the effect of silane flow rate and total gas pressure. For example, at a fixed R, the 

individual SiH4 and H2 flow rates can be changed so that one condition has a higher silane flow rate than 

the other. The films deposited at the same R were not the same and depended on the actual gas flow rates. 

This systematic study of the correlation between material structure and deposition conditions can help us 

understand the growth mechanism and adjust the growth conditions of nc-Si:H films. 

V.2. Sample and Experimental 

We studied 125 intrinsic Si films by HW-CVD on Corning 1737 25 × 25 mm2

 glass. All of the films were 

grown at the same substrate temperature, Ts = 240o

C. There were four groups with various silane flow 

rates of 3, 8, 16, and 22 sccm, respectively. The hydrogen flow rate for each group was varied from 0 to 

216 sccm. The total gas pressure was controlled by adjusting the throttling valve position (TVP) at 45%, 

50%, 60%, 70%, and 80%. The higher percentage leads to lower pressure. For the films with a low silane 

flow rate of 3 sccm, the thickness was controlled in the range of 1500-2000 Å, and the growth rate was in 

the range of 1-3 Å/s; for the films with silane flow rates of 8 and 16 sccm, the thickness was controlled in 

the range of 5000-6000 Å, and the growth rate was in the range of 5-13 Å/s; for the films with a silane 

flow rate of 22 sccm, the thickness was controlled in the range of 0.9-1.1 μm, and the growth rate was in 

the range of 8-21 Å/s. Two films were made as thin as the penetration depth of the 514.5-nm laser, 60-70 

nm, and these were used for calibrating the Raman spectra from back excitation. 

32

Raman spectra were measured using an XY triple spectrograph equipped with a LN-cooled CCD detector. 

The data were taken at room temperature using the 514.5-nm line of an argon-ion laser. The spectral 

resolution was approximately 0.5 cm-1 in a frequency range of 100 to 1000 cm-1. The frequency of the 

Raman lines was calibrated using the TO mode of a (111) c-Si wafer at 520 cm-1. The power of the laser 

was kept below the level that would thermally induce crystallization. The penetration depth of the 514.5-

nm light is ~60 nm for a-Si:H and is larger for nc-Si:H. Since the Raman spectra reveal the structure of 

the thin layer, measurements were also done from back-excitation for most samples. To depress the 

scattering light from the substrate, the light spot was focused on the film using a microscope. The c-Si 

grain size ranged from 7 to 30 nm, which was obtained from XRD; the crystalline volume fraction, Xc, 

was deduced using equations (5.3) and (5.4) with 0.4 < y(L) < 0.85. 

V.3. Results and discussion 

In general, the structure change is a function of the ratio of hydrogen-to-silane flows, R. Fig. 5.1(a-p) 

shows the crystalline silicon volume fraction, Xc, as a function of R for four groups of films at SiH4 flow 

rates of 3, 8, 16, and 22 sccm and with the total gas pressure varied with the throttling valve position 

(TVP) of 45%, 50%, 60%, 70%, and 80%. The solid and dotted lines indicate the Xc obtained from the 

top and bottom layer, respectively. The Xc is found from zero to >90% (including Xg.b. from zero to 

~30%) in the top layers. In agreement with the previous works, there is a threshold R for the structural 

transition from a- to nc-Si:H. The common features of the curves in Fig. 5.1(a-p) are: (a) Narrow structure 

transition region (∆R~1) from Xc= 0 to ≥50%. We will analyze this transition region in detail shortly. (b) 

The threshold R of crystallinity increases with increasing SiH4 flow rates. For instance, the threshold is 

R≤2 for low SiH4 flow rates of 3 and 8 sccm, as shown in Figs. 5.1(a-g), but no crystallinity has yet 

occurred at R=2 for high SiH4 flow rates of 16 and 22 sccm as shown in Figs. 5.1(i-p). (c) The Xc changes 

with the total pressure in the chamber in the following way: the films easily become nc-Si at the higher 

pressure (lower TVP) and lower silane flow rates of 3 and 8 sccm as seen by comparing Figs. 5.1(a) to 

5.1(d) and 5.1(e) to 5.1(h); whereas, at high silane flow rates of 16 and 22 sccm, a lower pressure (higher 

TVP) yields a higher Xc in the film as seen by comparing Figs. 5.1(i) to 5.1(j) or 5.1(m) to 5.1(n). 

These surprising results suggest that there is a “growth zone” for the crystallite, and it is our intention to 

study many deposition parameters, and not just R, to examine this further. We used back excitation to 

study the early growth of the film. At the early stage of the growth, Xc is much lower than in the late 

growth of the films as shown by the dotted lines in Fig 5.1. Whereas, in Fig. 5.1(d) for the film with R = 

72 and SiH4 = 3 sccm, we observe roughly the same Xc value. This is because the film thickness of 70 nm 

is close to the penetration depth of the 514.5-nm laser. For all of the films in this study, the FWMHs of 

the a-Si network from the bottom layers are in the range of 75-83 cm-1. These values are always broader 

than the 67-80 cm-1 for the top layers. This implies that the a-Si:H network near the glass substrate is less 

ordered compared to the top surface layer. 

33

100 

80 

0 

20 

40 

60 

80 

100 

0 

SiH4=3 sccm, TVP=45%

c-Si volume fraction (%) 

(a) 

0 

TVP=50%

(b) 

0 

TVP=60%

(c) 

0 

TVP=70%

(d) 

20 40 60 80 20 40 60 80 20 40 60 80 20 40 60 80 

H-dilution Ratio H-dilution Ratio H-dilution Ratio H-dilution Ratio


SiH4=8 sccm,TVP=45% 

0 5 10 15 20 

(e) 

TVP=50% 

(f) 

c-Si valume fraction (%) c-Si Valume fraction (%) c-Si volume fraction (%)

TVP=60% 

(g) 

TVP=70% 

80 

60 

40 

20 

0 

(h) 

0 5 10 15 20 0 5 10 15 20 0 5 10 15 20 

TVP=60% TVP=70% TVP=80% 

(j)


0 5 10 15 20


SiH4=16 sccm, TVP=50% 

(i) 

80 

60 

40 

20 

0 

0 5 

(k) 

0 5 10 15 20 

(l) 

0 5 10 15 20 10 15 20 

TVP=60% TVP=70% TVP=80% 

60 

SiH4=22, TVP=50% 

40 

20 

(m) 

0 5 10 15 20 0 

(n) 

5 10 15 20 0 

(o) 

5 10 15 20 

(p) 

0 5 10 15 20 

H-dilution ratio, R H-dilution ratio, R H-dilution ratio, R H-dilution ratio, R 

Fig. 5.1. Xc as a function of R at SiH4 flow rates of 3, 8, 16 and 22 sccm. The solid and dotted lines 

indicate the Xc was obtained from the top or bottom layer of ≥ 60 nm, respectively. 

34 

0 
35

Figs. 5.2(a) and 5.2(b) show the Raman spectra and their fitting functions for the films with SiH4 flow of 

3 sccm, TVP=45%, and R=1 and 2, respectively. The fittings for the TA mode peaked at 145 cm-1 and the 

LA mode peaked at 330 cm-1, which are also shown. The signal between the TO and LA modes is 

attributed to the LO mode. Comparing Fig. 5.2(a) and 5.2(b), one can see that the intensity ratio of 

TA/TO decreased as the film was more crystallized. This agrees qualitatively with the theoretical 

calculations. More importantly, we obtain the structure information from the TO mode according to 

equations (5.1) and (5.2). Figs. 5.2(c) and 5.2(d) show the enlarged TO mode between 450 to 550 cm-1 for 

the same samples in Figs. 2(a) and 2(b). The bandwidth of the TO mode, FWHM, and their peak 

frequency, ωTO, as a function of H-dilution ratio are plotted in Figs. 5.3(a) and 5.3(b) for all the films in 

this study. Generally, the FWHM decreases and the ωTO increases in the narrow structural transition 

region and then reaches a saturation value. Notice that the films with a high SiH4 flow rate of 22 sccm 

(the big crosses) show the narrowest FWHM and lowest ωTO. It is perhaps related to column-like 

structures. The changes can be seen more clearly for the low growth rate films with a SiH4 flow rate of 3 

sccm which are represented by solid dots in both Fig. 5.3(a) and 5.3(b).

100 200 300 400 500 600

SiH4=3sccm, R=1, TVP=45%

Raman intensity (a.u.)

Raman Shift (cm-1)

(a)

100 200 300 400 500 600

SiH4=3sccm, R=1, TVP=45%

Raman intensity (a.u.)

Raman Shift (cm-1)

(a)

460 480 500 520 540

SiH4=3sccm, R=1, TVP=45%

Raman intensity (a.u.)

Raman Shift (cm-1)

(c)

460 480 500 520 540

SiH4=3sccm, R=2, TVP=45%

Raman intensity (a.u.)

Raman shift (cm-1)

(d)

Fig. 5.2. Raman shift and their fittings from the films with SiH4 flow of 3 sccm, TVP=45%, 

(a) R=1, and (b) R=2. And the enlarged TO mode of the data in 2(c) and 2(d). 
R=3 R=3 

SiH4=3 sccm 

SiH4=8 

SiH4=16 

SiH4=22 

(a) 

85 482

(b)

FWHM of a-Si TO mode (cm-1) 

80 

75 

70 

65 

60 

a-Si peak frequency, 

ωTO (cm-1) 480 

478 

476 

474 

472 

55 470 

0 20 40 60 80 0 20 40 60 80 

H-dilution ratio H-dilution Ratio 

Figs. 5.3.(a) FWHM and (b) ωa-Si of the TO mode as function of R. The solid dots, dashes, open 

circles, and crosses represent the films with a SiH4 flow rate of 3, 8, 16, and 22 sccm, 

respectively. There are two regimes separated at R≈3: the fast changing regime when R <3 and 

the saturated regime when R>3. 

The improvement of SRO of the a-Si:H network in the a-to-nc-Si transition region is clearly shown in 

both the changes of FWHM and ωTO. In Fig. 5.3(a), the FWHM decreases from 80 to 67 cm-1 when 0 ≤ R 

≤ 3, and then remains at 67 cm-1 when R > 3. This indicates that the spread in the mean bond angle ∆Θ 

decreased from 10.8° to 8.6° according to equation (5.1). The corresponding frequency shift ωTO for the 

same films in Fig. 5.3(b) is from 475 cm-1 to 480 cm-1, which indicates that the ∆Θ decreased from 12.2° 

to 10.2° according to Eq. (5.2). The absolute values of ∆Θ are relatively large from the calculation, but 

the 2° improvement of the bond angle spread is consistent with the results from the decrease of the band 

width. The improvement of SRO of the a-Si network by increasing H-dilution is consistent with the 

observations of a narrowing of the valence band tail in our previous work. For most of the high-SiH4-

flow-rate films, indicated by the open circles and big crosses, the FWHM = 67 cm-1 does not change with 

R, but the ωTO spreads from 475 to 480 cm-1 at 0<R<3. Whereas, the FWHM is as narrow as 58.5 cm-1 

and the peak frequency as low as 472 cm-1 for the films deposited with the highest growth rate of ~20 Å/s 

(represented by the big crosses). The frequencies of the c-Si and g.b. peaks for those films are also low 

(see Figs. 5.4 and 5.5). Since a column-type growth occurs in such conditions, it implies that both the ωTO 

and the FWHM are sensitive to the column structures. To our knowledge, there are no calculations of how 

the phonon density of states changes with such structures. For the c-Si component, we generally found 

that the width of the c-Si TO mode for the nc-Si:H films is broader than that for the c-Si wafer, i.e., 8-10 

cm-1 compared to 4 cm-1. There were two exceptions for films with SiH4=3 sccm with TPV=70% and 

R=1.2 and 3. The reason is unclear why the FWHM was as narrow as 4 cm-1 for those two films. 

However, the peak frequency of the c-Si TO mode, ωc-Si, changes when R increases. For the films with 

SiH4 flow rate of 3 and 8 sccm, we found a blue shift from 516 to 520 cm-1 in the transition region, as 

shown in Fig. 5.4. This indicates that the crystallite size has increased from ~7 nm to ≥ 10 nm. For the 

films with a high SiH4 flow rate of 22 sccm, the big crosses indicate that the ωc-Si blue shift occurs from 

513 to 518 cm-1. Again, the relatively low frequency of the c-Si TO mode could be due to the columnlike structures. 

36

516 

518 

520 

522 

SiH4=3 sccm

 = 

=16 sccm

 =22 sccm 

515 

8 sccm

c-Si peak frequency, ω c-Si (cm-1) 514 

512 

0 20 40 60 80 

H-dilution Ratio, R 

495 

0 20 40 60 80 

SiH4=3 sccm 

SiH4=8 

SiH4=16 

SiH4=22 

H-dilution Ratio, R

g.b. peak frequency, ω g.b. (cm-1) 

510 

505 

500 

 Fig. 5.4. ωc-Si as a function of R. Fig. 5.5. ωg.b. as a function of R. 

Finally, for all the nc-Si:H films, the Raman TO mode always contains an intermediate component in 

addition to the a-Si and c-Si components. The intermediate component was fitted well using a Gaussian 

function with FWHM of 33 cm-1 for all the nc-Si:H films. The peak frequency of the intermediate 

component, ωg.b, as a function of R is plotted in Fig. 5.5. One finds that the value spreads from 500 to 510 

cm-1. For the low-growth-rate films, there is a blue shift of the ωg.b from 500 to 507 cm-1 in the transition 

region and then spreads between 505 to 507 cm-1 when R >3. Regardless of whether the intermediate 

component originates from the grain boundaries, small size (<3 nm) crystallites, or median-range ordered 

objects, the blue shift of the ωg.b can be attributed to the growth of the c-Si grain size. One may argue that 

Raman modes reflect SRO, and the phonon density of states is sensitive to the local bond angles. 

However, the scattering intensity is related to a coupling-parameter-weighted phonon density of states. 

Small differences between peak positions could be attributed to changes in median- or long-range 

interactions that are consequences of the coupling parameter. We could not clarify the origin of this 

intermediate component in this work. Whatever the origin is, one commonly takes the third component 

into account as part of the c-Si volume fraction as seen in equation (5.4). As shown above, the peak 

frequencies of all the three components of the TO modes show blue shift with an increase of H-dilution in 

the structural transition region. The compressive stress could be higher in films with higher hydrogen 

dilution ratios (the film easily peels off from the substrate), and the blue shift of the phonon bands could 

be partially related to an increase of the film stress. 

V.4. Summary 

Raman scattering studies were carried out on the transition films from amorphous-to-nanocrystalline 

silicon prepared by HW-CVD with varied silane and hydrogen flow rates and the T-valve position. The 

measurements were done for both front and back excitation. The TO mode was deconvoluted into a broad 

Gaussian function for the a-Si:H and a narrow Lorenzian function for the c-Si; a Gaussian function 

peaked at between 500 to 510 cm-1 with FWHM of 33 cm-1 for the intermediate component. We believe 

that the origin of the intermediate component is the grain boundaries and the small-size (< 3 nm) 

crystallites, and attribute it into the total c-Si volume fraction. The c-Si volume fraction depends not only 

on the H-dilution ratio, R= H2/SiH4, and the film thickness, but also on the SiH4 flow rate and total gas 

37

pressure which can be explained using the silane-depletion model. The detailed calculation of gas 

depletion will be reported in a future work. 

The most interesting finding is a narrow structural transition region of ∆R~1, in which the short-range 

order of the a-Si:H network is improved: the width of the TO mode decreased from 80 to 67 cm-1 and the 

peak frequency, ωTO, blue shift from 475 to 480 cm-1. These indicate that the bond-angle variation, ∆Θ, 

decreases from ~10 ̊ to ~8 ̊ when the crystallites start to grow. Furthermore, the crystallite grain size 

increases from ~7 nm to ≥ 10 nm in the transition region, evidenced by a blue shift of the c-Si mode, ωc-Si, 

from 516 to 520 cm-1. When R continually increases beyond the transition region, the SRO of the a-Si:H 

network, the c-Si volume fraction, and the c-Si peak frequency are somewhat saturated. However, the 

structure would be continually modified at a higher hydrogen dilution ratio. Other techniques such as 

TEM will assist in clarifying the further structural changes. 

Finally, when the films were deposited with a high growth rate ~20 Å/s using a high SiH4 flow rate, the 

FWHM of the TO mode for a-Si network can be as narrow as 58.5 cm-1 and the peak frequency as low as 

472 cm-1. A column-type growth occurs in such conditions, which implies that both the ωTO and the 

FWHM are sensitive to the column structures. Calculations are needed to find how the phonon density of 

states changes with such structures. 

V.5 Conclusion 

There is a narrow structural transition region of ∆R~1, in which the short-range order of the a-Si:H 

network is improved that, i.e., the bond-angle variation, ∆Θ, decreases from ~10 ̊ to ~8 ̊ and the 

crystallites start to grow. The crystallite grain size increases from ~7 nm to ≥ 10 nm in the transition 

region. 

VI . Micro-Raman Studies of Mixed-Phase Hydrogenated Silicon Solar Cells [14] 

VI.1. Introduction 

A significant light-induced enhancement in the open-circuit voltage (Voc) of mixed-phase silicon solar 

cells was observed. Since the average grain size is in the order of a few nano-meters in the mixed-phase 

material, the magnitude of the variation in Voc (∆Voc) depends on the intrinsic (i) layer thickness, the 

deposition temperature, the initial Voc, and the light soaking intensity. Under intense light soaking, a ∆Voc 

as large as 150 mV has been observed. Subsequent annealing of the solar cells substantially restored the 

original Voc values. In-situ PL spectral studies found that the PL intensity and peak-energy position 

associated with the amorphous component of the mixed-phase material increase upon light soaking. We 

proposed that a reduction of Xc or size is responsible for the Vo enhancement. However, recent studies by 

conventional Raman and X-ray diffraction spectroscopy found no observable changes before and after 

light soaking, prompting us to carry out further investigations. 

Raman spectroscopy is a sensitive tool that provides valuable structural information about a-Si:H 

materials. For the mixed-phase material, the Xc is low, and the distribution of the grains might be nonuniform. Using conventional Raman to detect the existence of nano-grains is difficult. However, microRaman is an effective tool for detecting nano-grains and providing their distribution information. In this 

paper, we present the micro-Raman date on the mixed-phase solar cells, which correlated well the 

structural non-uniformity to the physical model of light-induced increase of Voc observed in the mixedphase solar cells. 

38

VI.2 Sample and Experimental 

Single-junction n i p solar cells were deposited by United Solar onto 4 cm × 4 cm stainless steel (ss) 

substrates using a conventional rf glow discharge technique. Indium-tin-oxide (ITO) dotes of area 0.25 

cm

2 were deposited on the p layer as the top contact. The deposition parameters were adjusted to arrive at 

the amorphous-to-nanocrystalline transition region . 

The Raman spectra were collected under ambient conditions using the 514.5 nm line of an argon-ion 

laser. The penetration depth of the 514.5-nm light is ~60 nm for a-Si:H and is larger for nc-Si:H. The 

spectral resolution was approximately 0.5 cm-1 in a frequency range of 100 to 1000 cm-1. The frequency 

of the Raman lines were calibrated using the TO mode of a (111) c-Si wafer at 520 cm-1. A low laser 

power was used to avoid possibility of laser-induced crystallization. The incident beam was 

approximately three microns in diameter. At first, four Raman scans were taken on the bottom quadrant of 

the dots. The locations of the scans are shown in Fig. 6.1(a). Since non-uniformity appears to be more 

pronounced for the cells with Voc ~ 0.75 V, twenty scans were made in steps of 500 μm in the same 

bottom quarter of sample 13678-42, as indicated in Fig. 6.1(b). The laser beam was moved to a new spot 

on the ITO dot. This allowed us to scan the samples under the same initial conditions for each scan and 

thereby avoid any residual effects (such as light induced metastable effects) of laser radiation at the 

previously irradiated spot. 

a b 

cd 

1 2 3 4 5 

10 9 8 7 6 

11 12 13 14 15 

20 19 18 17 16 

(a) (b) 

Fig. 6.1. Schematics of (a) 4 scans for all samples, and (b) 20 scans for 13678-42. 

VI.3 Results and Discussion 

cm

A cell's Voc can be conveniently used to determine if it is in the amorphous or mixed phase. Cells with 

high, low, and medium Voc, corresponding to a-Si:H, nc-Si:H and mixed-phase Si solar cells were studied. 

Table VI.1 lists the Voc values and the Raman features for four scans on each cell. The four scans for 

sample 13687-32 with VOC = 0.981 V are the same and exhibit the typical amorphous signature, as seen in 

Fig. 6.2. The spectrum can be decomposed into four broad Gaussian distributions: one centered at 480 

-1 with a FWHM of 67 cm-1 representing the TO mode, one at ~ 140 cm-1 with FWHM of 100 cm-1 

representing the TA mode, one around ~330 cm-1 with FWHM of 117 cm-1 representing the LA mode, and 

one around ~ 440 cm-1 with FWHM of 100 cm-1 representing the LO mode. The four components in the 

spectrum are typical for high quality a-Si:H material. Figure 6.3 shows the Raman TO mode for sample 

13678-13 with a VOC= 0.674 V. All Raman spectra indicate a nc-Si:H material with a sharp c-Si TO mode 

around 516 cm-1 appears. We decompose the Raman spectrum to calculate Xc using Eq. (5.4) and obtain a 

volume fraction of 35%-56%. 

39

40

Table VI.1. Voc of the mix-phase solar cells and their Raman results 

Sample ID Voc (V) Raman lineshape for four scans 

13687-32 0.981 All broad a-Si peak at ~ 480 cm-1

13678-42 0.767 All broad a-Si peak at ~ 480 cm-1

13678-43 0.748 One with c-Si peak at ~ 516 cm-1

13678-13 0.674 All sharp c-Si peak at~ 516 cm-1

When the mixed-phase cells with VOC ~ 0.767 V and ~ 0.748 V were scanned by Raman, interesting 

results were observed as shown in Figs. 6.4 (a) and 6.4 (b). The four scans in Fig. 6.4(a) show typical 

amorphous features. In Fig. 6.4(b), three scans show the amorphous signature, but one scan clearly 

exhibits a sharp Lorentzian peak around 516 cm-1. The fact that only one of the eight spectra shows 

nanocrystallinity indicates that the material is very non-uniform in the micrometer scale. One may 

question if the cell in Fig. 6.4 (a) also contains nc-Si, because the VOC ~ 0.767 V corresponds to a mixedphase structure. The sample in Fig. 4 (a) was further scanned in 20 systematic steps of 500 μm illustrated 

in Figure 6.1 (b). Since the beam size is only 3 μm in diameter, there was no possible overlap between 

adjacent scans. 

 

100 200 300 400 500 600 700

Raman Intensity (a.u.)

Raman shift (cm-1) 

350 400 450 500 550 600

Raman shift (cm-1)

Fig. 6.2 Raman spectrum of 13687-32. Fig. 6.3 Raman spectrum of 13678-13. 

250 300 350 400 450 500 550 600 Raman Intensity (a.u.)

Raman shift (cm-1)

(a)

 

250 300 350 400 450 500 550 600 Raman Intensity (a.u.)

Raman Shift (cm-1)

(b)

Figure 6.4. Raman shifts of mixed-phase cells with (a) VOC = 0.767 V and (b) VOC = 0.748 V.
41

Figure 6.5. Raman Shift of 20 scans on a mixed phase cell 13678-42. 

The 20 Raman spectra shown in Fig. 6.5 exhibit typical Gaussian lineshapes centered at~480 cm-1. 

Several spectra show a hint of a nanocrystalline shoulder around 512 cm-1. Of the entire set of spectra, 

only one spectrum shows definite nanocrystallinity. This confirms the structural non-uniformity for both 

the mixed-phase cells with VOC ~ 0.767 V and ~ 0.748 V. For the c-Si peak on 13678-43 and 13678-42, 

the frequency was at 516.7 cm-1 and 514 cm-1, respectively. It is clear that the small c-Si grains are not 

uniformly distributed in the a-Si:H matrix but are clustered in a micrometer-size spot. In the two-diode 

model, a mixed-phase solar cell is equivalent to two diodes connected in parallel. One diode has the 

characteristics of an a-Si:H cell with Voc around 1 V, and the other one has the characteristics of an ncSi:H cell with a Voc below 0.5 V. It is assumed in the model that a certain area of the cell contains more 

nanocrystallites than the others. In this case, the lateral transport between the a-Si:H phase and the ncSi:H phase becomes negligible. The non-uniform distribution of nanocrystallites obtained by microRaman measurements in this study provides a direct supporting evidence for the assumption.

VI.4 Summary 

We used micro-Raman spectroscopy to study the degree and dimension of the structural uniformity of 

mixed-phase silicon solar cells. Low-laser power was used to avoid possible thermal effects. For a cell 

with Voc = 0.981 V, four scans on different locations of the cell show typical broad Gaussian lineshapes 

around 480 cm-1, indicative of an amorphous silicon structure. All scans of the cell with Voc = 0.674 V 

show a sharp Lorentzian peak around 516 cm-1, with an estimated 40% nanocrystalline volume fraction. 

For mixed-phase cells with Voc = 0.748 and 0.767 V, one of twenty and one of four spectra, respectively, 

350 450 550

1

10

11

20

350 450 550

2

9

12

19

350 450 550

3

8

13

18

350 450 550

4

7

14

17

350 450 550

5

6

15

16
exhibited both a broad Gaussian lineshape around 480 cm-1 and a sharp peak around 516 cm-1. This 

suggests that a large fraction of a-Si:H component but small fraction of nc-Si:H grains within the mm dot 

of the mixed phase cells. The spatial distribution of the nc-Si:H grains is clustered in a micrometer spot. 

The light-induced Voc enhancement has been observed in such mixed-phase cells. The structural nonuniformity observed in this micro-Raman study provides direct supporting evidence for the two-diode 

model recently proposed for explaining the light-induced Voc enhancement in mixed-phase solar cells. 

VI.5. Conclusion 

A small fraction of c-Si grains clustered within the mm dot in the mixed phase cells. This micro-Raman 

study provides direct evidence of the structural non-uniformity that causes the light-induced Voc 

enhancement in mixed-phase solar cells. 

VII. Toward understanding the degradation without light soaking in hot-wire a-Si:H thin films and 

solar cells [15] 

VII.1 Introduction 

The advantages of HW-CVD are the simplicity in gas decomposition, the absence of ion bombardment, 

and the high growth rate compared with the standard PE-CVD. High performance a-Si:H based solar cells 

and thin-film transistors have been made using HW-CVD. In addition, several new physical phenomena 

were observed in the HW thin films. For instance, the NMR studies showed a new hydrogen distribution 

in HW films compared to PE-CVD films, the internal fraction studies showed HW films have a much 

lower friction value than glass and are very close to the value of c-Si. In the high growth-rate films, the 

findings of the confinement effect on spin-spin interactions in nanogas systems and nanovoid-related 

large red shift of PL peak energy have been reported. The features of light soaking effects in HW a-Si:H 

films were distinguished from those in conventional PE-CVD films. The PC decreased only slightly and 

the dark conductivity increased in the films deposited at a substrate temperature ≥ 320 oC. The increase of 

the density of the metastable dangling bonds (∆Do) was found to be the same as that in PE-CVD films. On 

the other hand, it has been reported that in addition to the light-induced degradation the HW solar cells 

showed a few percent decrease in efficiency without exposure to the light. Most degradation in fill factor 

has shown in the air but not in the vacuum or in liquid nitrogen environment. Those results suggest a 

possible post-oxidation in the devices. The degradation without light soaking is a more serious problem in 

the HW-CVD samples than that in the PE-CVD samples. In this work, we studied in detail the variations 

in electronic properties of intrinsic HW films in terms of gas adsorption and thermal annealing in order to 

understand the phenomena of degradation in the dark. 

VII.2 Experimental Details 

Intrinsic HW a-Si:H films with low initial dark conductivity and large Ea ≥ 0.85 eV were selected as listed 

in Table VII.1. One standard PE-CVD film, A7119, was used for comparison. Group 2 was a pair of films 

with and without a thin SiNx passivition layer for post-oxidation study. Group 3 was one sample that was 

cut into three pieces for gas adsorption and thermal annealing studies. Corning-1737-glass substrate was 

used for all the samples. The coplanar metal contacts were first deposited on the glass. The details of the 

HW-CVD system are published elsewhere [12]. 

42

Table VI1.1. Sample information 


Group 

1 

Sample ID 

T773 

Ts (°C) 

360 

Growth Rate (Å/s) 

10 

T779 360 10 

 T837 400 12 

2 H1441 360 14 

 H1442 360 14 

3 H1540a 320 12 

H1540b 320 12 

H1540c

A7119a

 320 

200 

12 

1 a

 PE-CVD sample 

The measurements were carried out in the following experiments: (a) Dark conductivity measurement at 

300 K in environments of: vacuum (10-4 Torr), dry-oxygen gas, and air. Thermal annealing cycles were 

followed after storage in the above environments. (b) Conductivity activation-energy measurement with 

sample temperature ranging from 300 to 420 K. Before taken data, the sample was heated to 420 K for 30 

min to clean up the film surface adsorption. Then, the Ea was deduced from the equation σd = σoexp(-

Ea/kT). (c) Hot-probe measurement is a simple way to distinguish between n- and p-type semiconductors 

using a hot-probe such as a heated soldering iron, a cold-probe and a standard multimeter. Based on the 

thermal-electric effect of semiconductors, when applying the probes to n- or p-type materials one obtains 

a current in opposite directions. Finally, PC temperature dependence was studied for two typical HW- and 

one PE-CVD film. 

VII.3 Result and discussion 

For group 1 HW samples, the PC decreased only 0.5 

slightly and the dark conductivity increased upon 

light soaking. We used hot probe to exam the 0.0 

conductive type of the intrinsic samples. We found 

Current (10-14 A) 

that high Ts HW samples appeared slightly p-type 

initially and changed to n-type after defect 

-0.5 

creation. Figure 7.1 shows the results of the hot- -1.0 

probe measurement for sample T837. One can see 

that the sample at the initial state was slightly p- -1.5 

type, and then changed to slightly n-type after the 

metastable defect creation. In corresponding to its -2.0 

Ea value, the Fermi level moved up from 0.92 eV A B2 B6 

to 0.82 eV below the conduction band edge. History of the sample 

Fig. 7.1 The hot-probe data of sample T837. 

The Ea of T837 was as large as 0.92 eV. After the metastable defect creation at State B2 (light soaked 

using 200 mW/cm2

 white light for 2 hrs) and B6 (light soaked using 200 mW/cm2 white light for 6 hrs), 

the current became negative, and the Ea decreased to 0.82 eV. This change is in an opposite direction 

from standard PE-CVD film (A7119) in that the Ea increased from 0.73 eV to 0.77 eV after the 

metastable defect creation. The changes in the activation energy and the PC were partially recovered after 

° annealing at 160 C for an hour. The activation energy recovered to ~0.86 eV instead of 0.92 eV. So, there 

43

is some non-reversible degradation. In order to exam the effect of post-oxidation on the position of the 

Fermi level, we studied a pair of film with and without the SiNx passivition layer. We assumed that a thin 

layer of SiNx would prevent adsorption from air. The dark current as a function of temperature and 

annealing time are shown in Fig. 7.2. The 10 times difference in the value of the current between H1441 

and H1442 can be attributed to the interface states between SiNx and a-Si:H. This interface effect 

dominates the conductivity measurement on H1441. For solar cells, we have reported a 9.8% initial 

efficiency that became 9.2% after a few weeks of being stored in the dark. The degradation mainly came 

from the fill factor (FF). In order to check whether the degradation took place in the intrinsic layer, we cut 

sample H1540 into three pieces of H1540a, H1540b and H1541c and studied the current changes in the 

vacuum (x), pure oxygen (Ο) and the air (∆), simultaneously. Figures 7.3(a) and 7.3(b) show the changes 

of the dark current at 300 K and the corresponding activation energy, respectively. The crosses in Fig. 

7.3(a) show that the conductivity in the vacuum of 10-4 torr did not change much up to a month. However, 

the conductivity increased about an order of magnitude after following thermal cycles heated up to 420 K. 

The change in conductivity corresponded to Ea changes from 0.81 to 0.78 eV as shown in Fig. 7.3(b). The 

open circles show interesting changes of the conductivity in pure O2: the initial dark current has the same 

value as in the vacuum but increased and then decreased. Whereas, the current value at 300 K increased 

about an order of magnitude after temperature increased to 420 K; again the changes in conductivity 

corresponded to Ea changes as shown in Fig. 7.3(b). Finally, the open triangles show the conductivity 

measurement in the air: the dark current increased about 60 times more than that in the vacuum within 

one day and continuously increased in the air for 54 days. The current decreased an order of magnitude 

after Ea measurements in the air. We believe that this is due to de-adsorption of humidity at elevated 

temperatures. 

H1441.vacuum 

10-6 

10-7 

O2 

H1442.vacuum 

O2 

N2 

(a) 

Current (A)

H1441.vacuum 

H1441.O2 H1442.vacuum 

H1442.O2 

(b) 

Current (A)


10-7 

10-8 

10-9 

10-8 

10-9 

10-10 

10-11 

10-12 

2.4 2.6 2.8 3.0 3.2 3.4 0 10 20 30 40 50 60 70 

1000/T (1/K) Annealing time (min.) 

Fig. 7.2. (a) Current-temperature dependence and (b) current as a function of annealing time at 420 

K, for samples of H1441 and H1442 with- and without SiNx passivition layer. The samples were 

mounted in the vacuum and in pure oxygen or nitrogen, respectively. 

44

Current (10-12 A) 

102 

101 

100 

Vacuum O2 Air 

After 420 K

(a) 

0.66 

0.70 

0.74 

0.78 

0.82 

0 2 4 6 8 

Vacuum O2 Air 

E a (eV) 

(b) 

Vacuum 

10 12 14 0 500 1000 1500 

Aging time (hr) Aging time (days) 

Fig. 7.3 Degradation experiments: (a) the dark current at 300 K for samples of H1540a in vacuum (×), 

H1540b in pure O2 (�), and H1540c in the air (∆). The arrows indicate the current after Ea 

measurements. (b) The Ea obtained from the three samples after the current measurements at 300 K for 

about a month. The arrow indicate the measurements were done in vacuum for H1540c. 

As shown in Figs 7.3(a) and 7.3(b), the final values of both the current at 300 K and their activation 

energy Ea tended to reach the same value. The above results suggest that H2O from the air has a stronger 

effect than the pure O2. The source of the degradation in the dark can be related to trace oxygen 

incorporation at 300 K in the air or at 420 K in a 10-4 torr vacuum (we can not exclude the fact that the 

structural relaxation may also play a role) that results in an upward movement of the Fermi-level position. 

Consequently, the hole-mobility-lifetime product decreases and then the FF decreases. 

10-11 

10-10 

10-9 

10-8 

10-7 

10-6 

State A 

State B 

Photoonductivity (Ohm-cm)-1 

(a) 

Dark conductivity 

10-11 

10-10 

10-9 

10-8 

10-7 

10-6 

State A 

State B 

Photoonductivity (Ohm-cm)-1 

(b) 

Dark conductivity 

2 4 6 8 10 12 2 4 6 8 10 12 

1000/T (1/K) 1000/T (1/K) 

Figure 7.4 PC temperature dependence for (a) HW film T779 and (b) PE-CVD film A7119. 

45 
All of the above results showed that the changes in activation energy, in other words, the Fermi-level 

position were crucial to understanding the degradation in the dark. Since the gap states were crucial for 

the Fermi-level position, we compared the gap states between HW- and PE-CVD films in Figure 7.4 

using PC temperature dependence. To show the opposite movement of the Fermi-level positions in these 

two films, the dark conductivity curves are also plotted in Figs 7.4(a) and 7.4(b). We found that there are 

two extra PC enhancement peaks at ~200 K and 250 K for HW film, and the ~200 K peak disappeared 

after light soaking. No matter what those states are, the disappearance of the deep traps is consistent with 

the upward movement of the Fermi level. One may also notice that the PC value does not decrease at ~ 

room temperature at State B for the HW film but it does for the PECVD film. 

VII.4 Conclusions 

For HW a-Si:H samples, we have observed the decreases in the FF of the solar cells and the dramatic 

changes in the dark conductivity for the i-layer regarding the various gases and thermal cycles. Three 

effects could be involved in the changes that are: surface adsorption, oxygen incorporation and thermal 

relaxation. We observed that the surface adsorption of H2O in the air resulted in a great rising of the dark 

conductivity as shown in Fig. 7.3(a). The value of activation energy was slightly decreased after 

temperature cycles in the vacuum. When the activation energy decreased from ~0.9 eV to ~0.8 eV, the 

material changed from slightly p-type to slightly n-type, as shown in Fig. 7.1. This resulted in a decrease 

of the mobility-lifetime product of holes, and then a decrease of FF. We suggest that the degradation of 

both materials and solar cells may share the same origin. With prolonged storage time in the air at 300 K 

or being heated to 420 K in a 10-4 torr vacuum, oxygen incorporation resulted in the Fermi-level position 

moving upward. Essentially, the differences in both the microstructures and the gap states (see Fig. 7.4) in 

HW- from PE-CVD samples could be in response to the degradation in the dark. 

VIII. Micro-Raman Studies for Electrical-induced Structural Changes in Metal/μ c-Si/Metal 

Structures Made by Hot-Wire CVD [16] 

VIII.1. Sample and Experiment 

Boron-doped μc-Si:H thin films of 30-50 nm thickness were deposited on Cr-coated glass substrate at 

160 °C by HW-CVD. A W wire heated to 2000 °C was used to decompose silane, H2 and B3H6 (3.1% in 

helium) precursor gases. The flow rates for those gases are 3, 45 and 3 sccm, respectively. The devices 

have photolitho-graphically-defined Ag or Al top contacts with the lateral size varying from 3 to 40 μm. 

The thickness of both the bottom (Cr) and the top (Ag or Al) contacts was ~ 80 nm. I-V characteristics 

were recorded during switching with a Keithley 2400 SourceMeter. Switching was stimulated by a 

current-ramp from 10 nA up to 50 mA between the top and bottom contacts with a total ramp time of 

about 35 seconds. Before switching, the resistance of the device depended inversely on the device area. 

After switching, the resistance decreased more than three orders of magnitude and became independent of 

the area. The surface temperature on the device during stressing and switching was probed. 

A Dilor Triple Spectrograph XY Raman Spectrometer (Jobin Won-Spex ISA. Inc.) was used to analyze 

the material structures. Raman spectra were collected under ambient conditions using a 514.5-nm line of 

an argon-ion laser with a beam size of 3 μm in diameter. The laser power was more than one order of 

magnitude lower than the threshold power of laser-induced crystallization. The Raman frequency was 

calibrated using the TO mode of a (111) c-Si wafer at 520 cm-1. Two arrays of 40-μm-diameter Cr/μcSi/Ag (and Cr/μc-Si/Al) devices were studied. The microstructure of the μc-Si layer was analyzed after 

the metal top contacts were removed. 

46

VIII.2. Results 

8.2.1. The I-V characteristics of M/a-Si/M and M/μc-Si/M structures 

0.0 

1.0 

2.0 

3.0 

10-6 10-5 10-4 10-3 10-2 

Voltage (V) 

Current (A) 

(a) 

A 

C 

0.0 

1.0 

2.0 

3.0 

10-8 10-7 10-6 10-5 10-4 10-3 10-2 10-1 

Voltage (V) 

Current (A) 

(b) 

A 

C 

B 

+ /

×

2 + 

μ

Fig. 8.1. Typical switching voltage vs. current traces in (a) a 80 μm-diameter Cr/p a-Si:H Ag and 

(b) a ~ 5 5 μm Cr/p c-Si/Ag devices. 

Figure 8.1 shows the comparison of the typical switching voltage-current characteristic using a current 

source for (a) a Cr/a-p+

Si:H/Ag device of 5x103 μm

2 with a ramping time of ~ 0.5 sec and (b) a Cr/p+ 

μcSi/Ag device of 25 μm2 with a ramping time of ~ 35 seconds. The voltage across the device is plotted as a 

function of applied current bias. In the M/a-p+

Si:H/M devices, as shown in Fig. 8.1(a), when the current 

bias increases to point A (~ 20 μA) the voltage across the device suddenly drops from ~2.3 V to ~0.03 V 

at point C. Then the current near linearly depends on voltage. The switching I-V characteristic in Fig. 8.1a 

indicates the M/a-Si:H barrier is thermally brokendown from a high- to a low-resistance state. After the 

thermal breakdown, there is no longer a M/a-Si:H Schottky barrier but a M/M resistance that results in the 

I-V linear dependence. However, in the Cr/p+ 

μc-Si/Ag devices as shown in Fig. 8.1(b), at low bias (≤ 1 

μA), the voltage is almost proportional to the current. At intermediate biases (1 ~ 10 μA), the voltage 

changes nearly logarithmically with the current. As the bias is raised to about 20 μA, the voltage becomes 

saturated at about 2.2 V (point A). The voltage remains almost constant until the bias reaches about 20 

mA (point B) when the voltage drops to a lower value of about 1.0 V (point C). At the same time, the 

device resistance permanently decreased to about 500 Ω. The status of the device depends on the applied 

bias current. At bias ≤10 μA, the voltage-current characteristic is reversible indicating no switching 

processes yet. When the bias increases to 10 μA or a little beyond, hysterics is observed as shown by the 

dotted line in Fig. 8.1(b). This hysteretic behavior is repeatable, and little irreversible change occurs 

provided that the bias is never increased beyond the point A. When the bias current continually increases 

from 20 μA to 20 mA, the voltage saturated between A and B. An irreversible resistance decrease was 

observed at a point between A and B which implies irreversible changes of the structure. Finally, between 

points B and C, the voltage drops to 1.0 V with a resistance dropping to a low value. Whatever, this 

switching behavior is very different from what is observed in the a-Si:H devices where no saturation in 

voltage occurs before switching. And the voltage shows an instantaneous big drop once the bias reaches a 

critical value compared to a gradual and relatively small drop in the μc-Si:H devices. After switching, the 

resistance is considerably lower than that in a comparable μc-Si:H device. 

47

48

8.2.2. Raman studies before and after switching of the M/μc-Si/M structures 

Previous work on M/p+

a-Si:H/M devices was unclear as to whether there were any structural changes of 

the a-Si:H matrix during the forming process. In order to explore whether structural change took place 

during the switching in M/p+

μc-Si:H/M devices, we measured Raman spectra before and after electrical 

switching. Raman is a sensitive tool for measuring the degree of crystallinity in μc-Si:H materials. Two 

arrays of 40-μm-diameter Cr/p+

μc-Si:H/Ag and Cr/p+μc-Si:H/Al devices with a μc-Si:H layer of 30-50 

nm thickness were studied. The microstructure of the μc-Si:H layer was analyzed after the metal top 

contacts were removed. For the first array (Ag contact), the Raman spectra of seven non-switched and 

eight switched devices were measured. The Raman laser beam excited the center of each device. The 

Raman signals were weak compared to the signals from other μc-Si samples that have been studied. Long 

time accumulations (4000 sec for one spectrum) were needed to achieve an acceptable signal to noise 

ratio. Figure 8.2 shows the curve fitting of the average spectrum from (a) seven non-switched devices, (b) 

eight switched devices and (c) the overall average of the Raman spectra before and after switching. The 

data fitting results indicate that an average c-Si volume fraction increased from 51.4% to 54.2%, and the 

c-Si bandwidth became narrower from 12 to ~ 11 cm-1 after switching. 

The surface temperature on a device during switching was probed using a fine thermocouple that touched 

the top contact. It was seen that the surface temperature rises linearly in time from 25 °C to 40 °C during 

the period of voltage saturation. Finally, temperature drops dramatically with the voltage when the 

switching finished. The small device size and the likelihood of heat loss through the thermocouple 

preclude quantitative conclusions from the surface temperature measurements. The temperature in the 

bulk must be higher than at the surface. These observations suggest that the voltage saturation and the 

irreversible switching were closely related to sample heating. 

VIII.3. Summary 

P-type doped microcrystalline Si film deposited on Cr-coated corning 1737 glass by HW-CVD at a 

substrate temperature of 160 °C was used in metal/p+

μc-Si/metal switching devices. Ag or Al was used for 

the top contact with a diameter of 3-40 μm. After switching, the resistance decreased from >1 MΩ to <1 kΩ

and became independent of the area. Compared to M/p+

a-Si:H/M devices we found that (a) the switching in 

460 480 500 520 540

Raman shift (cm-1)

(a) f

c

 = 51.4%

460 480 500 520 540

0

20

40

60

80

100

Raman Shift (cm

-1)

(b) f

c

= 54.2%

460 480 500 520 540

OFF

ON

Raman shift (cm-1)

(c)

Fig. 8.2. The Raman spectra fitted using three functions for (a) the seven non-switched, (b) the 

eight switched devices, and (c) comparing the spectra in (a) and (b).
the μc-Si devices generally required lower applied voltage than comparable a-Si:H switches, (b) the I-V 

characteristic of the switching processes showed a saturated voltage region. We used Micro-Raman to study 

structural changes due to switching processes. The incident laser beam was 3 μm in diameter. Numbers of 

locations (10-25 spots) were scanned across each device. We found statistically that the average c-Si 

volume fraction increased from ~51.4 % to ~54%. We have described a conducting percolation path 

through the g.b. regions that connect the top and bottom contacts, which results in a low resistance state 

permanently. 

VIII.4. Conclusion 

By means of micro-Raman spectroscopy, we found a small but statistically reproducible increase in 

crystalline volume fraction after switching in metal/p+ 

μc-Si/metal devices. We conclude that an increase 

of crystallinity of the bottom of the i-layer was involved in the electrical switching processes. 

IX. Buffer-layer Effect on Mixed-Phase Cells Studied by Raman and PL Spectroscopy [17] 

IX.1. Introduction 

A light induced open-circuit voltage (Voc) increase has been observed in mixed-phase hydrogenated silicon 

solar cells, where the intrinsic layer contains a small amount of nanocrystalline grains in the amorphous 

matrix. An increase of PL intensity and a blue shift of the peak energy at 80 K have been observed 

previously in an in-situ light soaking experiment. The Voc increase was initially explained by light-induced 

structural changes, i.e. an expansion of the amorphous phase and a reduction of the crystalline phase. 

However, the structural change model explains the light-induced enhancement of volume on the order of 

0.1-0.01%, while the observed Voc increase is as large or greater than 15%. Both XRD and Raman did not 

find any observable light-induced structural changes in the mixed-phase materials. Meanwhile, microRaman found non-uniform distribution of the nanocrystalline grains on a micrometer-scale in the mixedphase cell. Experimental results also showed that the Voc enhancement can be suppressed by applying a 

reverse bias during light soaking, which suggested that the effect is driven by the recombination of photogenerated carriers, the same as the usual SWE. Furthermore, when a fully a-Si buffer layer (~ 500 Å) was 

inserted between the mixed-phase intrinsic and the p layers, the solar cells showed high initial Voc and no 

light-induced enhancement in Voc. Based on those observations, we have proposed a two-diode equivalentcircuit model, which explains most of the experimental results. A mixed-phase cell was considered as two 

diodes connected in parallel: one cell with an amorphous signature of Voc ~1.0 V and the other with a 

microcrystalline signature of Voc ~0.5 V. The lateral transport between the two cells is negligible. 

According to the experimental I-V characteristics for good and poor quality μc-Si cells, the Voc increase in 

a mixed-phase cell can be understood as the metastable defect creation in both μc-Si:H and a-Si:H diodes. 

Indeed, the increase of Voc is always accompanied with degradation of the fill factor (FF) and short circuit 

current density (Jsc). A thin a-Si:H (100 Å) buffer layer did not change the light-induced Voc increase, 

which was explained by carrier tunneling through the thin buffer layer. However, a thick (500 Å) buffer 

layer eliminated the light-induced Voc increase, since it blocks the current path through the microcrystalline 

phase. We use micro-Raman and PL spectroscopy to further study the effects of an a-Si:H buffer layer on 

the microstructure and electronic density of states in mixed-phase solar cells. 

IX.2. Sample and Experimental 

Mixed-phase silicon n-i-p solar cells were deposited using a RF glow discharge on specular stainless steel 

(ss) substrates. The i layer is about 250-nm thick. The top contacts are Indium-Tin-Oxide (ITO) dots with 

an active area of 0.25 cm2

. The solar cells in the first group were from sample L13869, where differences 

in the crystalline volume fraction and cell performance are due to a nonuniform deposition . In this sample, 

49

no a-Si:H buffer layer was deposited between the mixed-phase intrinsic layer and the doped layers. The 

cells in the second (sample L13868) and third (sample L13870) groups were deposited under the same 

conditions as in the first group except an a-Si:H buffer layer with a thickness of 100 Å and 500 Å, 

respectively, was deposited between the p layer and the i layer. The initial cell performances are listed in 

Table IX.1. In the initial state, the 100 Å thick buffer layer blocks the shunt current and increases Voc and 

FF for the mixed-phase solar cells (cell 11 and cell 13), but does not show an effect on the cell with a fully 

amorphous signature (cell 23). Increasing the thickness of the buffer layer to 500 Å, the initial Voc is still 

increased for the mixed-phase cells, but the FF decreases due to a kink in the J-V curve caused by 

inadequate tunneling through the a-Si:H buffer layer. After light soaking, the light-induced Voc increase 

was observed in both the cells with no buffer layer and with a 100 Å thick buffer layer, but was not 

observed in the cells with a 500 Å thick buffer. Details can be found elsewhere. 

Table IX.1. The initial performance of the mixed-phase solar cells with different a-Si:H buffer layers 

Sample # Cell # Buffer 

(Å) 

Pmax 

(mW/cm2

) 

Jsc 

(mA/cm2

) 

Voc 

(V) 

FF 

L13869 11 0 4.39 10.8 0.695 0.587 

13 0 5.47 10.3 0.852 0.621 

22 0 7.05 10.3 1.001 0.682 

L13868 11 100 4.77 10.9 0.725 0.604 

13 100 6.08 10.4 0.905 0.644 

23 100 7.08 10.4 1.003 0.680 

L13870 11 500 4.19 11.4 0.911 0.403 

13 500 6.83 11.3 0.970 0.625 

22 500 7.35 11.3 0.979 0.664 

cm

Raman measurements were made under ambient conditions on these solar cells using the 514.5-nm line 

from an argon-ion laser. The frequency was calibrated using the TO mode of a (111) c-Si wafer at 520 

-1. The penetration depth of the 514.5-nm light is ~ 600 Å for a-Si:H. Therefore, the structure of the top 

layer, which includes the buffer layer, was measured. Four spots were studied for each cell to evaluate the 

structural non-uniformity. PL spectra were measured in a temperature range of 80-300 K using a 632.8-

nm laser excitation. The penetration depth of the 632.8-nm light is ~1 μm for a-Si:H. Hence, the overall 

electronic states in the a-Si:H buffer layer and the mixed-phase intrinsic layer are probed. 

IX.3. Raman results 

V

Figure 9.1 shows the Raman spectra for the nine cells on the three samples. The three curves from the 

bottom to top in each plot are arranged with increasing microcrystallinity (decreasing Voc). One can see 

that the corresponding curves in Fig. 9.1(b) for the samples with a 100 Å thick a-Si:H buffer layer are 

almost identical to those in Fig. 9.1(a) for the samples with no buffer layer. This is because the 514.5-nm 

laser light penetrates through the 100 Å thick a-Si:H buffer layer, and the signature of the microcrystalline 

grains still appears on certain areas of the cell. However, no microcrystalline signature is observed on the 

spectra of the sample with a 500 Å thick a-Si:H buffer layer since the Raman scattering is mainly from 

the a-Si:H buffer layer. Concerning the relation between the Raman spectrum and Voc, we find that for 

oc ≥ 0.911 V, regardless of the crystallinity and the buffer layer, the TO mode peaked at ~ 480 cm-1 with 

a width of 68 cm-1 but without the peak or shoulder at 520 cm-1, indicating a typical amorphous structure. 

In this case, the crystallinity in the probed region near the i/p interface is below the measurement 

sensitivity. 

50

51

Figure 1. Raman spectra for the cells on (a) L13869, (b) L13868, and (c) L13870, corresponding to no 

buffer layer, 100 Å thick a-Si:H buffer layer, and 500 Å thick a-Si:H buffer layer, respectively. 

IX.4. PL results: 

9.4.1. PL spectra at low temperature: We now show the effect of crystallinity in the mixed-phase i-layer 

on the PL lineshape. Figure 9.2 shows the PL spectra at 80 K for the three cells on sample L13869. As 

300 350 400 450 500 550 600 650

L13869-11

L13869-13

L13869-22

Raman Shift (cm-1)

300 350 400 450 500 550 600 650

L13868-11

L13868-13

L13868-23

Raman Shift (cm-1)

300 350 400 450 500 550 600 650

L13870-11

L13870-13

L13870-22

Raman Shift (cm-1)

Fig. 9.1. Raman spectra for the cells on (a) L13869, (b) L13868, and (c) L13870, corresponding to no 

buffer layer, 100 Å thick a-Si:H buffer layer, and 500 Å thick a-Si:H buffer layer, respectively. 

0.8 1.0 1.2 1.4

L13869-11 L13869-13 L13869-22

PL relative intensity at 80 K (a.u.)

Photon Energy (eV)

(a)

0.8 1.0 1.2 1.4

(b)

0.8 1.0 1.2 1.4

L13869-11

L13868-11

L13870-11

Photon Energy (eV)

(a)

0.8 1.0 1.2 1.4

(b)

Figure 9.2. PL spectra at 80 K for cells 11, 13, Figure 9.3. PL spectra of the low Voc cells of 

and 22 on sample 13869 for (a) linear and L13869-11, L13868-11, and L13870-11 for 

(b) semi-log plots. (a) linear and (b) semi-log plots.
the crystallinity increases, the main band becomes narrower; and the low energy (<1.2 eV) signal 

increases as seen clearly in the semi-log plot. When a portion of the microcrystalline phase increases, the 

band tail of the a-Si:H matrix becomes narrower, and extra radiative recombination in the g.b. region 

appears. We then show the effect of the a-Si:H i/p buffer layer on the PL lineshape. Figure 9.3 plots the 

PL spectra of cell 11 from the three samples. These cells are located at the same position on the substrate, 

but with different thickness of the a-Si:H buffer layer (0, 100 Å, and 500Å). The a-Si:H buffer layer 

makes the PL main band broad, and the low energy PL band from the g.b. regions is suppressed, which is 

similar to decreasing the crystallinity as shown in Fig. 9.2. 

I

9.4.2. PL intensity temperature dependence: The differences in the PL features among the nine cells can 

also be seen from the temperature dependence of the PL intensity (IPL). As shown in Fig. 9.4 for mixedphase cells, the PL spectrum generally contains three components: a main band from a-Si:H at high 

energy with a strong temperature dependence (Ο), a g.b. (∆) band at low energy with a weak temperature 

dependence, and a defect band (�) that appears only at T>180 K. For the cells with no a-Si:H buffer layer 

(sample L13869), with the increase of crystallinity (from cell 22 to 11), the PL intensity of the g.b. band 

increases, and the slope of the IPL-T plot of the a-Si:H component becomes steeper. By fitting the data to 

PL = Io exp(-T/TL), we obtained a PL characteristic temperature of TL~20 K for cell 11, which is smaller 

than those from the other two cells (~22 K). A smaller TL implies a narrower band tail. Therefore, the PL 

intensity result is consistent with the PL lineshape, where cell 11 has the narrowest bandwidth as shown 

in Fig. 9.2. For the cells with a 100 Å thick a-Si:H buffer layer (sample L13868) as shown in Fig. 9.4 (b), 

the a-Si:H main-band and defect-band for the three cells are close to each and the g.b. band still appears 

PL contribution from the a-Si:H buffer layer. The spectra in Fig. 9.4(c) show typical characteristics for 

intrinsic a-Si:H films. Only the a-Si:H main-band and the defect-band are observed. The g.b. band is in 

the same order as in Fig. 9.4(a) but with a relatively weak intensity, which is due to the completely 

suppressed due to the PL contribution from the 500 Å thick a-Si:H buffer layer. In addition, the PL 

spectra for these three cells show the common features of the a-Si:H, including the TL around 20-22K, the 

PL peak energy around 1.40-1.42 eV at 80 K, and the temperature coefficient of the peak energy around 

1.5-2.0 meV/K. These data indicate a good quality a-Si:H matrix for the cells 

50 100 150 200 250 300 

L13869-11 

L13869-13 

L13869-22 

Integral PL intensity, IPL (a.u.) 

T (K) 

(a) 

100 150 200 250 300 

L13868-11 

L13868-13 

L13868-23 

T (K) 

(b) 

100 150 200 250 300 

L13870-11 

L13870-13 

L13870-22 

T (K) 

(c) 

50 50 

Fig. 9.4. PL intensities vs. temperature for cells of (a) L13869, (b) L13868, and (c) L13870, where circles 

represent the a-Si:H main band, triangles the g.b. band, and squares the defect band. 

52

9.4.3. in-situ light soaking effect


0.8 1.0 1.2 1.4 1.6 

13869-11(A) 

13869-11(B) 

Relative PL intensity (a.u.) 

Photon Energy (eV) 

0.18 eV(A), 

0.20 eV(B) 

(a) 

0.8 1.0 1.2 1.4 1.6 

L13868-11(A) 

L13868-11(B) 

Photon Energy (eV) 

0.24 eV(A), 

0.24 eV(B) 

(b) 

0.8 1.0 1.2 1.4 1.6 

L13870-11(A) 

L13870-11(B) 

Photon Energy (eV) 

FWHM 

=0.29 eV 

(c) 

Fig. 9.5 PL spectra at 80 K before and after light soaking for cells of (a) L13869-11, (b) L13868-11, 

and (c) L13870-11. 

V

After the PL was measured in the initial state (state A), in-situ light soaking was done at room 

temperature for 4 hours using a 200 mW/cm2

 white light with an infrared cut-off filter. The PL spectrum 

was then measured at temperatures from 80 K to 300 K in the light soaked state (state B). Correlated with 

oc increase an increase of PL intensity at 80 K have been observed in an in-situ light soaking 

experiments in mixed-phase solar cells. It was explained by light-induced structural changes, i.e. an 

expansion of the amorphous phase and a reduction of the crystalline phase. Figure 9.5 shows the PL 

spectra at 80 K before and after light soaking for cells L13869-11, L13868-11, and L13870-11. For the 

cell without a-Si:H buffer layer, we find that the PL main band becomes slightly wider and the PL total 

intensity also increases slightly in Fig. 9.5(a), consistent with our previous observation. For cell L13868-

11 in Fig. 9.5(b), no detectable change is found on the PL spectra, which could be due to the opposite 

effect of the thin a-Si:H buffer layer (PL intensity decreases upon light soaking) and mixed-phase 

intrinsic layer (PL intensity increases upon light soaking) on PL. When the a-Si:H buffer layer was as 

thick as 500 Å, the light soaking effects in the buffer layer dominate. The cell L13870-11 in Fig. 9.5(c) 

shows a large change in both the lineshape and intensity. The total PL intensity is decreased by ~40% and 

the relative intensity of the defect band increases, which is normally observed in a-Si:H. 

IX. 5. Summary 

We use micro-Raman and PL spectroscopy to study the effects of an a-Si:H buffer layer at the i/p 

interface of the mixed-phase silicon solar cells. We find that the signature of the crystalline 520 cm-1 

mode still appears on the Raman spectrum for the cells with a 100 Å thick a-Si:H buffer layer; but it 

completely disappears for cells with a 500 Å thick a-Si:H buffer layer. At 80 K, the PL spectral lineshape 

reflects the features of the electronic states in the band tails. The characteristics of the PL spectra of the 

mixed-phase cells are a narrower main band than the standard a-Si:H band and an extra low energy band 

from the grain boundary region. As the thickness of the a-Si:H buffer layer increases, the PL main band 

becomes broader, and the low energy band is depressed. We find that, after light soaking, the PL main 

band is slightly broadened for the cells with no a-Si:H buffer layer, almost no change for the cells with a 

100 Å thick buffer layer, and a remarkable decrease in total PL intensity for the cells with a 500 Å thick 

53

buffer layer. In addition, the PL intensity of the defect band increases after light soaking for the cells with 

a 500 Å thick buffer layer, where light-induced defect generation in the a-Si:H buffer layer masks the 

changes in the mixed-phase intrinsic layer. 

IX. 6. Conclusion: The Raman and PL results could explain previous observations of the effect of an thin 

a-Si:H buffer layer on the performance and metastability against light soaking for mixed-phase solar cells. 

X. Correlation of Hydrogenated Nanocrystalline Silicon Microstructure and Solar Cell 

Performance[18] 

X.1. Introduction 

a-Si:H solar panels have been used for both terrestrial and space applications. One limitation of such 

panels is the light-induced degradation caused by the metastability of a-Si:H. Although much effort has 

been made to optimize solar cell structures, such as using a spectrum splitting triple-junction structure, the 

light-induced degradation still exists in a-Si:H based solar cells. In order to make high efficiency and 

high stable thin film solar cell, hydrogenated nanocrystalline silicon (nc-Si:H) solar cell has received a 

great deal of attention in the last decade due to its lower light-induced degradation and higher current 

density than a-Si:H cell. Initial efficiencies of over 10% for a single-junction nc-Si:H solar cell and over 

14% for a-Si:H/nc-Si:H double-junction solar cell have been reported. At United Solar, several 

techniques have been used to increase the deposition rate and improve cell performance. It is well 

accepted that a high quality material is essential for achieving high efficiency solar cells. Therefore, it is 

very important to find a direct correlation between material properties and device performance for 

improving cell efficiency. In previous studies, Raman and PL spectroscopy have been used to characterize 

the microstructure and electronic states in nc-Si:H materials and solar cells. The Raman TO mode was 

used to deduce the crystalline volume fraction (Xc). For the PL spectra of nc-Si:H, a relatively narrow 

band at low energy was found to be the characteristic peak of nc-Si:H. To distinguish this low energy 

peak from the defect peak in conventional a-Si:H materials, we measured the temperature dependence of 

PL from 15 K to 300 K [19] and found that both the intensity and the peak energy position of the low 

energy PL peak in nc-Si:H are higher at low temperatures than at high temperatures. This behavior is in 

contrast to the defect peak in a-Si:H where the defect PL has a weak temperature dependence and is 

undetectable at low temperatures. In nc-Si:H, when the temperature was increased from 15 to 180 K, the 

intensity of the nanocrystalline related peak was rapidly quenched, and the peak energy shifted from 1.0 

eV to 0.83 eV. Above 180 K, the defect peak appeared and overlapped with the PL from grain boundary 

regions. The behavior of the low energy PL peak in nc-Si:H was explained in terms of the carrier 

thermalization in an exponential band tail in the grain boundary region. The tail width kT0 was found to 

be 20-24 meV, which is about half the value in typical a-Si:H. We used PL and Raman to study the 

relationship between the material properties and the device performance of nc-Si:H solar cells made using 

radio frequency (RF), modified very high frequency (MVHF), and microwave (μW) glow discharge. In 

addition, not only 632.8-nm but also 442-nm laser lines are used for PL excitation to distinguish the 

difference in electronic density of states between the top layer and bulk. 

X.2 Sample 

Several n-i-p nc-Si:H solar cells were deposited on Ag/ZnO back reflector using various techniques at 

different deposition rates. Indium-tin-oxide ITO dots with an active area of 0.25 cm2

 were deposited on 

the p layer as the top contact. J-V characteristics was measured under an AM1.5 solar simulator at 25 °C. 

Quantum efficiency was measured from 300 nm to 1000 nm to calibrate the Jsc. Table X.1 summarizes 

the solar cell deposition conditions and J-V characteristics. Raman measurements were made directly on 

54

the solar cells under the ambient condition using the 514.5-nm line from an argon-ion laser. The 

penetration depth of the 514.5-nm light is ~60 nm for a-Si:H and is larger for nc-Si:H. 

Table X.I. Preparation conditions and solar cell performance of the nc-Si:H solar cells studied. 

Dep. 

method 

Sample 

ID# 

d 

(μm) 

Rate 

(Å/s) 

Eff 

(%) 

Jsc 

(mA/cm2

) 

Voc 

(V) 

FF Comments 

RF 13917 1.1 1 6.66 23.43 0.472 0.593 Flat H2 dilution, ambient 

stable 

14085 1.0 1 7.00 22.55 0.493 0.630 H2 dilution profiling, 

ambient stable 

VHF 11667 1.1 3 6.85 23.73 0.464 0.622 

Flat H2 dilution, thick i/p 

buffer, ambient stable 

10973 0.9 3 4.92 22.39 0.415 0.530 Flat H2 dilution, ambient 

stable 

10976 1.1 3 5.00 23.26 0.406 0.530 Flat H2 dilution, ambient 

unstable 

11538 1.1 3 4.55 24.92 0.406 0.450 Flat H2 dilution, ambient 

stable 

11486 0.7 12 3.20 17.05 0.381 0.493 Flat H2 dilution, ambient 

unstable 

μW 7073 0.8 20 3.03 15.22 0.390 0.510 Flat H2 dilution, ambient 

stable 

X.3. Raman and PL results 

Table X.2. Summary of the Raman and PL analyses of the nc-Si:H solar cells. 

Dep. Sample Xc Xg.b IPLc/(IPLc+IPLa) TL (K), 633 nm. TL (K), 442 nm 

method ID# (%) (%) 633 nm 442 nm nc-peak. a-peak nc-peak. a-peak 

RF 13917 80.6 35.8 68.7 92.8 19.3 17.4. 17.8 -

14085 53.4 13.1 32.5 3.4 20.7 21.7 22.0 18.1 

11667 40.4 5.0 26.3 2.1 21.7 23.7 20.2 15.2 

VHF 10973 74.8 30.8 5.1 24.1 15.6 21.1 15.2 -

10976 79.0 34.5 32.9 87.0 19.9 23.3 18.7 -

11538 77.8 34.5 25.1 95.0 17.8 23.0 16.7 -

11486 72.2 31.1 1.0 18.5 31.2 21.5 22.2 -

μW 7073 37.6 14.5 4.0 3.7 36.9 22.8 29.1 -

I

where Xc, and Xg.b. are the total nanocrystalline volume fraction included in the grains and the grain 

boundaries and. the grain boundary volume fraction calculated from Raman spectra, respectively, 

PLc/(IPLc+IPLa) the fraction of the nanocrystalline PL intensity, and TL the characteristic temperature of 

PL as described in the text. 

Figure 10.1 shows the Raman spectra of the TO mode for the eight cells listed in Table X.1. The PL 

spectra for the same samples were measured in a temperature range of 80-300 K using 632.8-nm and 442-

nm laser line excitations to probe the electronic properties in the bulk of the nc-Si:H intrinsic layer and 

the top layer near the i/p interface, respectively. The PL spectra were plotted in Fig. 10.2 for comparison 

with the Raman results. Table X.2 lists the calculated Xc and Xg.b. values from Fig.10.1. 

55

56

As shown in Fig. 10.2, most of the nc-Si:H solar cells show two separated PL peaks at 80 K. The broad 

peak at ~1.4 eV originates from the tail-to-tail radiative recombination in the a-Si:H region, while the 

narrow peak at ≤ 0.9 eV is related to the nanocrystalline grains and from the tail-to-tail radiative 

recombination in the grain boundary regions. The relative PL intensity of the grain boundary region is 

represented by the fraction of IPLc/(IPLc+IPLa), where IPLc and IPla are the integrals of the low and high 

energy PL peaks, respectively. One finds that IPLc/(IPLc+IPLa) using 442-nm excitation change in the same 

350 400 450 500 550 600

13917

14085

11667

10973

10976

11538

11486

7073

Raman Shift (cm-1)

0.6 0.8 1 1.2 1.4 1.6

13917

14085

11667

10973

10976

11538

11486

7073

Photon energy (eV)

 Figure 10.1. Raman spectra of the eight nc-Si:H Figure 10.2. PL spectra of the same nc-Si:H solar 

 solar cells. cells at 80 K excited by 632.8-nm (solid) and 

442-nm (dashed) laser beams.
manner as Xg.b. does. Furthermore, the width of the electronic band tails, kTo, could be obtained from the 

temperature dependence of the corresponding PL intensities. According to the carrier thermalization 

model, kTo=kTLln(νoτr), where TL, νo, and τr are the slope of IPL-T plot, the attempt-to-escape frequency, 

and the carrier's recombination lifetime, respectively. The TL can be deduced by fitting the experimental 

data into IPL = Io exp(-T/TL). Table X.2 also summarizes the analysis of PL data for the eight nc-Si:H 

cells. 

X.4 Discussion 

We first compare the two nc-Si:H cells deposited using RF glow discharge. The main difference is that 

the sample 13917 was deposited using a constant hydrogen dilution, while the sample 14085 using a 

profiled hydrogen dilution with decreasing hydrogen dilution during the deposition. As shown in Table I, 

the hydrogen dilution technique has improved Voc and FF significantly, but there is a slight loss in the Jsc. 

A recent study using the hydrogen dilution profiling technique has resulted in an efficiency of 8.4% of the 

nc-Si:H single-junction cell and 13.4% of a-Si:H/nc-Si:H double-junction cell. Raman results indicate 

that sample 14085 has much lower crystalline volume fraction than sample 13917, which is consistent 

with the observed higher Voc and lower Jsc in the profiled cell than in the flat one. However, we need to 

keep in mind that the Raman measurement with 514.5-nm excitation only detects the top layer near the i/p 

interface. The crystalline volume fraction may be over-estimated for the sample with flat hydrogen 

dilution, but under-estimated for the sample with the profiled hydrogen. Another interesting observation 

is that the grain boundary volume fraction is relatively lower in the sample 14085 than in the sample 

13917, which implies that a low percentage of grain boundary is desirable for solar cell performance. 

Meanwhile, the hydrogen dilution profiling (#14085) results in a clear difference of the electronic density 

of states in the top layer from the bulk as shown in the PL spectra, i.e. the amorphous-related peak 

dominants and the nanocrystalline-related peak diminishes when the excitation was changed from the 

long wavelength (632.8-nm) to the short wavelength (442-nm). This is opposite to the samples with flat 

hydrogen dilution (#13917), where the nanocrystalline-related peak increases at the top layer. 

The MVHF deposited nc-Si:H cells (#11667) with a thick a-Si:H i/p buffer layer showed an improved cell 

performance due to reduced shunt current. Raman measurement showed a low crystalline volume 

fraction with very little of grain boundary component and PL showed a low nanocrystalline related peak 

at the i/p interface layer. Both results are consistent with the expected effect of the a-Si:H buffer layer. 

We found that some unoptimized nc-Si:H solar cells showed a significant ambient degradation related to 

post-oxidization. Cells 10973 and 10976 are made under two conditions such that one produces no 

ambient degradation and the other one does. Raman results show that sample 10976 has slightly high Xc, 

whereas, PL spectra are very different. Sample 10976 has a much higher nanocrystalline related PL peak 

than sample 10973, especially, at the top layer, indicating a high non-uniformity along the growth 

direction . It has been showed that a high nanocrystalline volume fraction with large grains easily causes 

micro-cracks, post-oxidization, and consequently the ambient degradation. 

Now we discuss the band tail width deduced from PL temperature dependence. As listed in table II, the TL 

related to the amorphous PL peak is in the range from 17 to 23 K for all the eight cells in the bulk (632-

nm excitation), and in the top layer (442-nm excitation) TL is ~ 17±2 K for the two high efficiency cells. 

The TL for the nanocrystalline related PL peak could be categorized into two types. In the cells with 

reasonable performance, the TL is smaller than 23 K and changes very little from the bulk to the top layer. 

In contrast, in the high rate MVHF (#11486) and the microwave cell (#7073) have large TL values (31-37 

K) in the bulk. Since the band tail width kTo is proportional to TL, the large TL in the poor nc-Si:H solar 

cells implies that a broad band tail in the nanocrystalline region leads to a poor cell performance. 

Normally, the band tail in the nanocrystalline region is considered from the grain boundary states. 

Therefore, finding ways to improve the grain boundary quality (less defects and less distorted bonds) 

would be beneficial for improving the nc-Si:H solar cell performance. 

57

X.5. Summary 

I

We used Raman and PL spectroscopy to study the relationship between the material properties and the 

solar cell performance of hydrogenated nc-Si:H. The crystalline volume fraction was deduced from the 

Raman spectrum. Generally, a high Xc leads to a high short circuit current density and a low open circuit 

voltage. PL spectra were measured using 632.8-nm and 442-nm laser lines. There are two distinguished 

PL peaks at 80 K, one at ~1.4 eV originating from the amorphous region, while the other at ≤ 0.9 eV from 

the nanocrystalline grain boundary regions. Generally, the intensity fraction of this low energy PL peak, 

PLc/(IPLa+IPLc), was larger for 442-nm than 632.8-nm excitation, indicating an increase in crystallinity 

along the growth direction. However, for the best initial performance cells obtained by H2 dilution 

profiling and the i/p buffer layer, the intensity fraction IPLc/(IPLa+IPLc) decreased from the bulk to the top 

i/p interface. The Raman and PL results give insight into the correlation between the microstructures and 

the cell performance, and verified that properly-controlled crystallinity in the intrinsic layer and buffer 

layer at the i/p interface layer are important for optimizing nc-Si:H solar cells. 

The Raman results show that a high Xc leads to a high Jsc and low Voc in nc-Si:H solar cells made with 

constant hydrogen dilution. Xc increases as film grows for flat hydrogen dilution. However, hydrogen 

dilution profiling, with dilution decreasing as film grows, lowers Xc as measured by 514-nm light with a 

penetration depth of 60 nm. The PL spectra excited by 632-nm and 442-nm laser beams indicate that the 

crystallinity increased along the growth direction for the flat hydrogen diluted samples. The top layers in 

the hydrogen dilution profiled nc-Si:H cells, however, have a higher amorphous PL signal than in the 

bulk. The improved cell efficiency by hydrogen dilution profiling indicates that a properly controlled Xc 

distribution along the growth direction is a very useful technique to obtain high efficiency. The band tail 

width in the grain boundary region was deduced from the PL intensity temperature dependence. The 

results show that a narrow band tail in the nanocrystalline phase correlates to improved cell efficiency. 

X.6. Conclusion 

The Raman and PL results give insight into the correlation between the microstructures and the cell 

performance, and verified that properly-controlled crystallinity in the intrinsic layer and buffer layer at the 

i/p interface layer are important for optimizing nc-Si:H solar cells. 

58

FURTHER EFFORTS

The P.Is. Daxing Han, Keda Wang have had response the LOS of Thin-Film PV Partnerships Program, 

No.RXL-4-44205 under a project title of Correlation between Thin Film Properties and their Solar Cell 

Performance Studied by PL and Raman Spectroscopies from Department of Physics, Boston College 

(BC) 

We address the category of Specialized Contributors in a title of Correlation between Thin Film 

Properties and their Solar Cell Performance Studied by PL and Raman Spectroscopies. A study of the 

Staebler-Wronski Effect in the intrinsic layer of solar cells in relation to solar cell degradation is included. 

We choose the PL and Micro-Raman spectroscopies as the main characterization methods to support thinfilm PV industry R&D in improvement of solar cell’s performance and reliability. PL and Raman are 

non-contact and non-destructive techniques that can be used on both thin film and solar cell structures. It 

allows us to find the correlation between the i-layer properties and the solar cell performance that is 

crucial for understanding the device physics and then improving cell design and optimizing preparation 

conditions. Our PL and Raman techniques are able to probe the uniformity in growth direction. In 

addition, Micro-Raman probes the horizontal uniformity. So far, we have studied a-Si and film-Si in 

collaboration with United Solar Ovonic Corporation and National Renewable Energy Laboratory. The 

same techniques can be used to study CdTe and CIS. 

If awarded a subcontract, BC will participate in National R&D Team activities and continue to cooperate 

with the team members. 

59

Publications 

1. Hydrogen dilution profiling for hydrogenated microcrystalline silicon solar cells, Baojie Yan,


Guozhen Yue, Jeffrey Yang, Subhendu Guha, D. L. Williamson, Daxing Han and Chun-Sheng 


Jiang, to be published in APL (2004).


2. Micro-Raman studies for electrical-induced structural changes in metal/μc-Si/metal structures made 

by hot-wire CVD, Brittany Huie, J.R. Weinberg-Wolf, E.C.T. Harley, Keda Wang and Daxing Han, 

Jian Hu, Pauls Stradins, Howard M. Branz, and Qi Wang. To be published in Thin Solid Films 

(2004). 

3. MICROSTRUCTURE EVOLUTION WITH THICKNESS AND HYDROGEN DILUTION 

PROFILE IN MICROCRYSTALLINE SILICON SOLAR CELLS, Baojie Yan, Guozhen Yue, 

Jeffrey Yang, Subhendu Guha, Keda Wang and Daxing Han, MRS symp. Proc. 808, A8.5 (2004), in 

press. 

4. Increase of crystallinity during electrical switching in metal/μc-Si/metal structures, Jian Hu, Pauls 

Stradins, Howard M. Branz, Qi Wang, J.R. Weinberg-Wolf, E.C.T. Harley, Chris Lawyer, Brittany 

Huie and Daxing Han, MRS symp. Proc. 808, A4.25 (2004), in press. 

5. Correlation of Hydrogenated Nanocrystalline Silicon Microstructure and Solar Cell Performance, 

Keda Wang, Anthony Canning, J.R. Weinberg-Wolf, E.C.T. Harley, Daxing Han, Baojie Yan, 

Guozhen Yue, Jeffrey Yang, and Subhendu Guha, MRS symp. Proc. 808, A9.53 (2004), in press. 

6. Buffer-layer Effect on Mixed-Phase Cells Studied by Micro-Raman and Photoluminescence 

Spectroscopy, Andrea Hilchey, Chris Lawyer, keda Wang, Daxing Han, Baojie Yan, Guozhen Yue, 

Jeffrey Yang, and Subhendu Guha, MRS symp. Proc. 808, A4.38 (2004), in press. 

7. High-throughput chemical vapor deposition system and thin-film silicon library, Qi Wang,


Fengzhen Liu, Daxing Han, Macromolecular Rapid Communications, Jan. 2004.


8. Raman study of thin films from amorphous-to-microcrystalline silicon prepared by hot-wire 

chemical vapor deposition, Daxing Han, J. D. Lorentzen, J. Weinberg-Wolf, L. E. McNeil, and Qi 

Wang, J. Appl. Phys., 94, 5 (2003) 2930-2936. 

9. Micro-Raman studies of Mixed-Phase Hydrogenated Silicon Solar Cells, J. Owens, D. Han, B. Yan, 

J. Yang, K. Lord, and S. Guha, Amorphous and Heterogeneous Silicon-Based Films, MRS 

Symposium Proc. 762, edited by J.R. Abelson, G. Ganguly, H. Matsumura, J. Robertson, and E. 

Schiff, 339 (2003). 

10. Toward understanding the degradation without light soaking in hot-wire a-Si:H thin films and solar 

cells, Qi Wang, Keda Wang, and Daxing Han, MRS Symposium Proc. 762, edited J.R. Abelson, G. 

Ganguly, H. Matsumura, J. Robertson, and E. Schiff, 399 (2003). 

11. Combinatorial approach to thin-film silicon materials and devices, Qi Wang, Leandro R. Tessler,


Helio Moutinho, Bobby To, John Perkins, Daxing Han, and Howard M. Branz, MRS Symposium


Proc. 762, edited by J.R. Abelson, G. Ganguly, H. Matsumura, J. Robertson, and E. Schiff, 413 


(2003).


12. Micro-Raman Measurements of Mixed-Phase Hydrogenated Silicon Solar Cells, J. Owens, D. Han, 

B. Yan, J. Yang, K. Lord, and S. Guha, Proc. NCPV review meeting (2003, 3, 21-24, Denver, CO). 

13. Correlation between cell performance and film properties as function of growth rate, G. Ganguly,


Daxing Han, et al, Proc. NCPV review meeting (2003, 3, 21-24, Denver, CO).


14. Narrow Gap a-SiGe:H Grown by HWCVD, Brent P. Nelson, Daxing Han, M. Boshta, B. Alavi, 

Rubin Braunstein, Don L. Williamson, and Yueqin Xu, Proc. NCPV review meeting (2003, 3, 21-24, 

Denver, CO). 

15. Hydrogen Distribution, Nanostructures and Optical Properties of High Deposition Rate Hot-Wire 

CVD a-Si:H, Jonathan Baugh, Keda Wang, Daxing Han, and Yue Wu, Thin Solid Films 430 (2003) 

95-99. 

16. Neutral Dangling Bonds May Not Be The Dominant Recombination Centers For Photoconductivity 

In Hot-Wire a-Si:H, Daxing Han, Guozhen Yue, Qi Wang, Tatsuo Shimizu, Thin Solid Films 430 

(2003) 141-144. 

60

17. Photo- and electro-luminiscence of a-Si:H and mixed-phase alloys, Daxing Han and Keda Wang, 

“Critical Review of Amorphous and Microcrystalline Silicon Materials and Solar Cells” a special 

issue of Solar Energy Materials and Solar Cells, 78 (2003) 181-233. 

18. Hydrogen structures and the optoelectronic properties in transition films from amorphous to 

microcrystalline silicon prepared by hot-wire chemical vapor deposition, Daxing Han, Keda Wang, 

Jessica M. Owens, Lynn Gedvilas, Brent Nelson, Hitoe Habuchi and Masako Tanaka, J. Appl. Phys. 

93 (2003) 3776-3783. 

19. Qi Wang, Eugene Iwaniczko, Jeffrey Yang, Kenneth Lord, Subhendu Guha, Keda Wang, and 

Daxing Han, "WIDE-GAP THIN FILM SI n-i-p SOLAR CELLS DEPOSITED BY HOT-WIRE 

CVD", 29th PVSC IEEE Orlando, USA, 2002. 

20. A two-domain model of light-induced structural changes in hydrogenated amorphous silicon, 

Jonathan Baugh and Daxing Han, Phys. Rev. B66 (2002) 115203. 

21. Optical Properties of i-layers as a Function of Growth Rate in Correlation to the Performance of aSi:H Solar Cells, Keda Wang, Jessica M. Owens, Jennifer Weinberg-Wolf, Daxing Han, Lynn 

Gedvilas, and Gautam Ganguly, Amorphous and Heterogeneous Silicon-Based Films, MRS 

Symposium Proc. 715, edited by J.D. Cohen, H. Matsumura, J.R. Abelson, and J. Robertson, pp 571-

576 (2002). 

22. Characterization of Microcrystalline Transition from Amorphous Silicon as a Function of Hydrogen 

Dilution and Substrate Temperature of Hot-wire CVD, Keda Wang, Haoyue Zhang, Jing Zhang, 

Jessica M. Owens, Jennifer Weinberg-Wolf, Daxing Han, Lynn Gedvilas and Brent Nelson, 

Amorphous and Heterogeneous Silicon-Based Films, MRS Symposium Proc. 715, edited by J.D. 

Cohen, H. Matsumura, J.R. Abelson, and J. Robertson, pp 147-152 (2002). 

23. Correlation of Light-induced Enhancement of Open-Circuit Voltage and Structural Change of 

Heterogeneous Silicon Solar Cells, Jeffrey Yang, Kenneth Lord, Baojie Yan, Arindam Banerjee, 

Subhendu Guha, Daxing Han and Keda Wang, Amorphous and Heterogeneous Silicon-Based Films, 

MRS Symposium Proc. 715, edited by J.D. Cohen, H. Matsumura, J.R. Abelson, and J. Robertson, 

pp 601-610 (2002). 

24. Nanovoid-related Large Red Shift of Photoluminescence Peak Energy in Hydrogenated Amorphous 

Silicon, Daxing Han, Guozhen Yue, Keda Wang, Jonathan Baugh, Yue Wu, Yueqin Xu, and Qi 

Wang, Appl.Phys. Lett. 80, 40-42 (2002). 

25. Effects of Hydrogen Dilution and i-layer Thickness on a-Si:H Solar Cells Studied by 

Photoluminescence Spectroscopy, Guozhen Yue, Daxing Han, Jeffrey Yang, Kenneth Lord, Boajie 

Yan, and Subhendu Guha, NCPV Program Review Meeting, (Oct, 2001, Denver, CO). 

26. Link of the Light-induced Structural Changes and the Conductivity Changes in Hydrogenated 

Amorphous Silicon --a Two-domain Model, Jonathan Baugh and Daxing Han, NCPV Program 

Review Meeting, (Oct, 2001, Denver, CO). 

27. Confinement effect on spin-spin interactions in nanogas systems, Jonathan Baugh, Alfred 

Kleinhammes, Daxing Han, Qi Wang & Yue Wu, Science, 294, issue 5546, 1505-1507 (2001). 

61

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62

ABSTRACT 

In high growth rate (≥ 50 Å/s) HW-CVD a-Si:H films, for the first time, we show gaseous molecules in 

nanovoids (~2% volume fraction of tube-like nanoscale voids), and demonstrate that confinement on the 

nanometer scale generates NMR effects that have never been observed in macroscopic systems. In the 

same system we found the PL peak red shift. We suggest that highly strained bonds on the inner surfaces 

of the nanoscale voids form broad conduction-band tail states that are responsible for the PL red shift. 

We characterized the structural transition from a- to nc-Si as function of H-dilution, thickness and Ts of 

both HW- and PE-CVD films using IR, Raman, PL, CPM/PDS and Ea et al. We found not only the c-Si 

volume fraction but also the g.b. and microstructures play an important role in the properties of the i-layer 

and their solar cell performance. We found a narrow structural transition zone in which the bond-angle 

variation, ∆Θ, decreases from 10 ̊ to 8 ̊. For nc-Si samples, we found a characteristic low energy PL peak 

and proved that is originated from the g.b. regions. 

V

Using micro-Raman, we found the structural non-uniformity in the mixed-phase solar cells that showed 

oc enhancement after light soaking. Using micro-Raman, we also found the slight increase of 

crystallinity in M/μc-Si/M devices after current forming. 

In cooperation with USOC, we studied the correlation between the microstructure, the DOS, and the solar 

cell performance before and after light soaking using Raman and PL with varied excitation wavelength 

and explained the microscopic origin of the optimized conditions for the solar cell preparations. The best 

cells have been prepared at near the edge of the transition from amorphous to microcrystallin silicon. 

More detailed studies of such cells can be found in our last final technical report (1998-2001). 

63

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Microscopic Mechanism of the Staebler-Wronski Effect in a-Si Films 

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1 October 2001–30 September 2004 

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D. Han 

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NREL/SR-520-37990 

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PVB55101 

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Department of Physics 

University of North Carolina 

Chapel Hill, North Carolina 

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14. ABSTRACT (Maximum 200 Words) 

In high growth rate (≥ 50 Å/s) HW-CVD a-Si:H films, for the first time, we show gaseous molecules in nanovoids 

(~2% volume fraction of tube-like nanoscale voids), and demonstrate that confinement on the nanometer scale 

generates NMR effects that have never been observed in macroscopic systems. In the same system we found the 

PL peak red shift. We suggest that highly strained bonds on the inner surfaces of the nanoscale voids form broad 

conduction-band tail states that are responsible for the PL red shift. We characterized the structural transition from ato nc-Si as function of H-dilution, thickness and Ts of both HW- and PE-CVD films using IR, Raman, PL, CPM/PDS 

and Ea et al. We found not only the c-Si volume fraction but also the g.b. and microstructures play an important role in 

the properties of the i-layer and their solar cell performance. We found a narrow structural transition zone in which 

the bond-angle variation, ∆Θ decreases from 10 ̊ to 8 ̊. For nc-Si samples, we found a characteristic low energy PL 

peak and proved that is originated from the g.b. regions. Using micro-Raman, we found the structural non-uniformity 

in the mixed-phase solar cells that showed Voc enhancement after light soaking. Using micro-Raman, we also found 

the slight increase of crystallinity in M/μc-Si/M devices after current forming. 

15. SUBJECT TERMS 

PV; Staebler-Wronski effect; amorphous silicon (a-Si:H); solar cells; microscopic mechanism; photoluminescence (PL); 

Raman spectroscopies; manufacturer; thin film; light-induced degradation; 

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