Physical, electrical, and optical properties of SF-PECVD-grown hydrogenated microcrystalline silicon...

Abstract Characterization results on hydrogenated

microcrystalline silicon (lc-Si:H) thin films grown in a

Saddle Field (SF) PECVD system are presented. The

microcrystalline content of the films is controlled by

the application of a positive electrical bias to the film

growth surface. The results of photoluminescence,

atomic force microscopy, infrared-absorption, and

electrical conductivity studies are presented. The

results correlate to the changing microcrystalline

content of the films in the same way as when micro-

crystalline content is influenced through growth

parameters such as hydrogen dilution in other CVD

techniques.

1 Introduction

The microcrystalline content of Saddle Field (SF)-

PECVD-grown lc-Si:H shows dependence on many

growth parameters (pressure, hydrogen dilution, flow-

rate, plasma power density), much as samples grown in

other CVD systems [1]. Previous work on the growth

of lc-Si:H in the SF-PECVD system has also shown

that the electrical potential at the substrate surface is

crucial to the formation of the microcrystalline phase.

The substrate surface must either have a positive

voltage applied to it [2] or be electrically floating [3] to

encourage the growth of microcrystalline material. The

technique of applying a voltage to the growth surface

only yields microcrystalline material under specific

pressure and voltage conditions (E. Johnson et al.,

submitted). It has been proposed that this effect is due

to the reduced ion bombardment of the growth surface.

The importance of ion-bombardment during lc-Si:H

growth has been studied in other growth technologies

[4], and elevated ion bombardment generally reduces

the microcrystalline content of the bulk and can even

result in a layered structure in the resulting films [5].

Energetic ions amorphize the top-most layer and the

transformation to microcrystalline material occurs in a

subsurface region. In this way, ion bombardment

affects not just the microscopic properties of the

material, but will affect devices through the creation of

undesired interfaces. Addressing this effect is crucial if

the SF-PECVD technique is to be applied to large area

electronic growth, and this work explores the effect on

material properties that result in attempting to mini-

mize ion bombardment through the use of an applied

electrical bias.

In this work, the optical and electrical character-

ization of matched sets of lc-Si:H films grown on test-

appropriate substrates is presented. The films are

grown in a SF-PECVD system using growth conditions

near the threshold of microcrystalline growth and the

surface potential is actively controlled with an external

voltage source. We present here the results of tem-

perature-dependent photoluminescence (PL), atomic

force microscopy (AFM), Fourier transform infrared

absorption (FTIR), and four-point temperature-

dependent conductivity. We correlate this data to

E. V. Johnson Æ S. Hoogland Æ E. Klem Æ N. Kherani ÆS. Zukotynski (&)Department of Electrical and Computer EngineeringUniversity of Toronto, Toronto, Ontario, CanadaM5S 3G4e-mail: [email protected]

Present address:E. V. JohnsonLPICM, Ecole Polytechnique, Palaiseau Cedex 91128,France

J Mater Sci: Mater Electron (2006) 17:789–799

DOI 10.1007/s10854-006-0025-0

123

Physical, electrical, and optical properties of SF-PECVD-grownhydrogenated microcrystalline silicon with growth surfaceelectrical bias

Erik V. Johnson Æ Sjoerd Hoogland ÆEthan Klem Æ Nazir Kherani Æ Stefan Zukotynski

Received: 9 January 2006 / Accepted: 11 March 2006� Springer Science+Business Media, LLC 2006

previously measured microcrystalline content (as

determined by Raman scattering). We show that the

use of substrate bias during the growth of microcrys-

talline material in a SF-PECVD system impacts all its

physical, optical, and electrical properties to a signifi-

cant extent, and that this impact can be correlated with

the microcrystalline phase transition.

2 Experimental procedure

All of the samples for this study were grown in a

SF-PECVD system, which employs a semi-transparent

central anode to extend the mean free path of electrons

and increase their likelihood of ionizing and dissoci-

ating collisions. Twenty sets of samples were grown for

this study. All of the samples were grown under similar

conditions, tabulated in Table 1.

Further detail about the growth conditions and the

topology for the application of a substrate surface bias

can be found in previous publications [2]. For these sets

of samples, only the chamber pressure and the sub-

strate surface electrical potential (Vsub) were actively

varied. The chamber pressure was varied from 100 to

300 mTorr and Vsub from 150 V to 250 V.

The PL spectra of the lc-Si:H samples deposited by

SF-PECVD on f-SiO2 were taken using a tunable

Spectra Physics Mai Tai diode-pumped Ti:Sapphire

laser operating at 800 nm, and frequency doubled with

a KTP crystal to 400 nm. The sample was cooled

through contact with a liquid He cooled cold-finger,

and the temperature was measured with thermocou-

ples both at the cold-finger and on the sample surface.

The temperature difference between the two was sig-

nificant, so the cold finger temperature is considered a

lower limit on the sample temperature and the surface

temperature an upper limit.

The surface morphology of the films deposited on

f-SiO2 were analyzed through AFM using a Digital

Instruments Nanoscope in tapping mode. The sampled

area of the film was 750 nm · 750 nm.

The films to be analyzed through FTIR were

deposited on c-Si wafers for sufficient transparency in

the infrared. The error introduced by coherent reflec-

tion in the film is ignored [6], so the absolute values of

the atomic density may be overestimated as the films

are very thin. However, the relative density is the more

important parameter in this study.

The films grown on f-SiO2 were deposited in the

Van Der Pauw configuration and were analyzed

through temperature dependent four-point conductiv-

ity. These measurements were taken over a tempera-

ture range from 390 K to 430 K—well below the

sample growth temperature but in the temperature

range for extended state conduction—and were done

in high vacuum at an estimated pressure of 10)6 Torr.

3 Results and discussion

3.1 Temperature-dependent PL

The different PL signatures of a-Si:H and lc-Si:H

clearly define the two materials. In a-Si:H, PL shows a

broad dominant peak at 1.3–1.5 eV which is attributed

to tail state recombination [7], and another much

weaker peak at 0.9–1.0 eV which may be attributed to

defect states at mid-gap [8]. In lc-Si:H, a narrow

0.9 eV peak dominates, particularly at low-tempera-

tures [9], and there is a second, much broader peak at

~1.2 eV. The low energy peak is attributed to a

recombination center, due to its narrow linewidth,

which indicates a crystalline or ordered character to

the system. Authors have alternatively declared this

peak a result of dangling bond luminescence in the

a-Si:H matrix, or to a transition between bandtail states

in the lc-Si:H gap [10]. Savchouk et al. [10] used poly-

Si deposited on glass by pyrolysis at 625�C. Post-

deposition hydrogenation increased the intensity of

the PL by a factor of 3–4, indicating the suppression of

a non-radiative recombination channel, most likely

defects at grain boundaries. Komuro et al. [11] have

demarcated three distinct peaks in the PL spectrum

of lc-Si:H: one below 0.8 eV, one between 0.9 eV

and 1.0 eV, and the tail-to-tail amorphous peak at

1.2–1.4 eV. They observed a much shorter lifetime for

the lower energy peaks, and maintain that the 0.8 eV

peak is due solely to the a-Si:H defects. Merdzhanova

et al. [12] have simultaneously observed both the

0.8 eV and the 0.9–1.0 eV peaks but attributed the

peak separation to interference effects.

The transition from one material phase to the other

over a series of depositions due to changing conditions

can be followed by observing the changing PL peaks,

as explored by Yue et al. [13] with samples grown by

HW-CVD with varying degrees of hydrogen dilution.

They have attributed certain luminescence peaks to

different types of boundaries—a boundary between

c-Si islands and the a-Si:H matrix at ~1.0 eV, and the

Table 1 Growth conditionsFlow rate (FR) = 5 sccmDilution ratio (R),SiH4:H2 = 1:10Substrate temperature(Tsub) � 180�CAnode current(IAN) = 20 mA

790 J Mater Sci: Mater Electron (2006) 17:789–799

123

other between c-Si grains at 0.9 eV. A closer exami-

nation of their data shows three peaks, as identified in

[11] and [14], where the peak at 0.9 eV is dominant at

high microcrystalline fractions and shifts towards

1.0 eV with increasing amorphous content. A second

peak emerges near 1.25 eV with decreasing micro-

crystalline content, and a third peak near 1.4 eV in

amorphous samples.

So in summary, a total of four distinct peaks can be

identified at the following locations:

(1) £0.8 eV attributed to dangling bond recombina-

tion in a-Si:H.

(2) 0.9–1.0 eV, attributed to either the electronic tail

recombination in lc-Si:H grains or recombination

at grain boundaries (0.9 eV for GB’s between

crystallites and 1.0 eV for GB’s between crystal-

lites and amorphous matrix).

(3) 1.2–1.3 eV, attributed to a shifted a-Si:H tail-

to-tail transition peak but only present when

some microcrystalline fraction occurs.

(4) 1.3–1.4 eV, attributed to tail-to-tail recombina-

tion in a-Si:H.

Han et al. [14] performed both front and back side

illumination of the films for PL and observed signifi-

cantly different PL in each case. In the same sample,

the 1.2–1.3 eV peak is dominant for film side illumi-

nation, while the ~1.4 eV peak is dominant for back-

side illumination. This suggests that the 1.2–1.3 eV

peak is indicative of either protocrystalline material, or

very small crystallites.

The coincidence of the PL peaks in a-Si:H

and lc-Si:H make data interpretation difficult, but the

emergence of the peak at 0.9–1.0 eV in lc-Si:H at

low-temperatures is a good diagnostic measure of the

predominance of lc-Si:H in the material. This peak

shows a strong temperature dependence, whereas the

a-Si:H peak is much more temperature stable [9].

In Figs. 1 and 2 are shown the results for tempera-

ture-dependent PL of two films grown on f-SiO2 in the

SF-PECVD system for this study. The sample of Fig. 1

was grown at P = 250 mTorr and Vsub = 200 V and has

a strong microcrystalline Raman character, whereas

the sample of Fig. 2 was grown at P = 250 mTorr and

Vsub = 175 V and has a strongly amorphous Raman

character.

The data presented have been corrected for the

system baseline and grating anomalies. It can be clearly

seen that the microcrystalline sample shows a charac-

teristic peak between 0.9 eV and 1 eV expected of a

microcrystalline sample. This peak, which decreases

with increasing temperature, is due to recombination at

the grain boundaries between crystalline domains. As

well, in these samples, a narrow, smaller shoulder can

be seen on the low energy side of the peaks at low

temperature. This peak corresponds to the peak at or

below 0.8 eV which has been explained as recombi-

nation through mid-gap defects in a-Si:H (dangling

bond to self-trapped hole state in valence band). The

fact that this peak is still present despite the lack of an

a-Si:H peak above 1.3 eV is an interesting phenome-

non, but difficult to comment on as this feature occurs

right at the edge of the range of the Ge detector.

Fig. 1 Temperaturedependent PL of sampleshowing microcrystallineRaman character, grown onf-SiO2 at 250 mTorr and200 V. Lines represent a bestfit for the peak above 0.9 eV(contribution of low-energypeak to fit not shown)

J Mater Sci: Mater Electron (2006) 17:789–799 791

123

In contrast to the PL spectrum of the lc-Si:H film,

the a-Si:H film shows a dominant peak characteristic of

bandtail-to-bandtail recombination in a-Si:H, as well as

a much smaller peak due to recombination at the

crystalline–amorphous grain boundaries just below

1.0 eV. A small peak due to recombination through

mid-gap defects can be seen near 0.8 eV.

A qualitative comparison of the two sets of data is

presented in Fig. 3. The data for the microcrystalline

sample has been magnified by a factor of 4, however

the alignment of the collection optics was modified

when the samples were changed so no good quantita-

tive comparison between the two sets of data can be

done, including any comments about PL efficiency.

The temperature dependence of the height of the PL

peaks of both films is presented in Fig. 4. The PL

efficiency of the films decreases with increasing tem-

perature, although it can be seen that the microcrys-

talline peak decreases much more dramatically than

the amorphous peak. This is consistent with results

from the literature on a-Si:H and lc-Si:H films [9].

The rapid quenching of the microcrystalline peak in

comparison to the amorphous peak promotes the

model suggesting recombination through tail-states

Fig. 3 Data for bothmicrocrystalline sample(peaks on the left, 200 V,200 mT) and amorphoussample (peaks on the right,175 V, 200 mT)

Fig. 2 Temperaturedependent PL of sampleshowing amorphous Ramancharacter. Sample grown onf-SiO2 at 250 mTorr and175 V. Lines shown are bestfit to peak above 1.1 eV(contribution to fit of peaknear 0.95 eV not shown)


123

in lc-Si:H tail states rather than through GB states.

The much shallower tail states of the microcrystallites

would be depopulated more quickly through thermal

activation of carriers than would be the much deeper

amorphous tails.

Additionally, the temperature dependence of the

location in energy of the PL peak is shown in Fig. 5. A

dramatic difference can be seen in the low temperature

peak locations for the sample grown at Vsub = 200 V

(showing microcrystalline character) and the one

grown at Vsub = 175 V (showing amorphous charac-

ter). The PL peak for the amorphous sample shows a

gradual red-shift with increasing temperature, whereas

the microcrystalline sample does not show as dramatic

a shift. It should be noted that the two highest data

points for the microcrystalline sample in Fig. 5 have a

much lower intensity (noted in Fig. 4) and may con-

stitute a separate peak entirely, rather than a shifted

peak.

The red-shift of the amorphous peak is attributed to

the depopulation of the shallow tail states through

thermal activation of carriers to the conduction and

valence bands [10]. This depopulation leaves behind

carriers at lower energies and thus results in a red-shift

of the spectrum. The lesser shift in the microcrystalline

sample, (along with the narrower line-width) indicates

Fig. 4 Temperaturedependence of peak PLintensity for sample showingmicrocrystalline characterand sample showingamorphous character

Fig. 5 PL peak location forsamples showing bothamorphous andmicrocrystalline character


123

a more ordered population is producing this PL peak,

which supports both the models using tail-state

recombination in the crystallites as well as GB

recombination.

The results of the temperature dependent PL study

on the films demonstrate the powerful effect that

substrate bias plays on the resulting film properties in

the SF-PECVD system. It must be restated that the

two films under study were grown under identical

growth conditions to similar thickness, except for a

difference in ion bombardment energy reduced by an

externally applied substrate bias. These films show PL

properties that encompass those exhibited by films

grown under dramatically different growth conditions

in other systems.

3.2 Atomic force microscopy

AFM is an important diagnostic tool for the growth

of lc-Si:H. It allows characterization of the surface

morphology and this gives information about the

microcrystalline content of the film as explored in the

work of Collins et al., who observed that the transition

to microcrystalline growth is accompanied by a

roughening of the growth surface [15, 16].

AFM characterization was performed on the sam-

ples grown on f-SiO2. In Fig. 6(a)–(f) are presented the

scans for these samples. There is clearly a correlation

between the surface roughness and the microcrystal-

linity in this limited set of samples. The two films that

showed high microcrystalline content (deposition con-

ditions of 250 mTorr/200 V (Fig. 6a) and 200 mTorr/

225 V (Fig. 6d)) also show much more surface rough-

ness. A plot of the roughness layer thickness depen-

dence on the microcrystalline content is presented in

Fig. 7.

Clearly, the microcrystalline content of the films and

the surface roughness are strongly correlated, as

described in the literature for lc-Si:H films grown in

other CVD systems. The microcrystalline content of

the SF-PECVD films has been solely modified through

the application of a substrate surface electrical bias.

3.3 FTIR spectroscopy—hydrogen content and

microcrystallinity in lc-Si:H films

The hydrogen content and the configuration of

hydrogen in lc-Si:H also gives interesting information

about the structure of the material. Kroll et al. [17]

studied the hydrogen concentration of films of varying

microcrystalline fraction (controlled by H2-dilution).

As hydrogen dilution of the supply gas is increased

and the films become protocrystalline, the hydrogen

content first peaks, and then decreases to a few atomic

percent. As well, the configuration of the remaining

hydrogen changes from mostly mono-hydride (Si–H)

configuration to di-hydride (Si–H2). This effect has

also been noted by Fujiwara et al. [18] at interfaces

between lc-Si:H and ZnO, indicating that it is intrinsic

to the phase transition of proto-Si:H into lc-Si:H.

Fig. 6 AFM scans of films grown on f-SiO2 with varyingpressures and substrate bias voltages. (a) 250 mTorr and200 V, (Xl c = 64.4%), (b) 250 mTorr and 175 V, (Xlc= 4.7%),(c) 200 mTorr and 250 V, (Xl c = 44.8%), (d) 200 mTorr and225 V, (Xl c = 63.0%), (e) 200 mTorr and 200 V, (Xlc = 14.5%),and (f) 200 mTorr and 175 V, (Xl c = 21.5%)


123

Silicon–hydrogen bonds contribute to the absorp-

tion spectrum through three modes: stretching (near

2000 cm)1), bending (near 850 cm)1) and wagging

(640 cm)1). Isolated Si–H bonds contribute to the

wagging peak at 640 cm)1, and also to a stretching

peak at 2000 cm)1. The presence of silicon dihydride is

indicated by a bending mode near 880 cm)1, while

silicon trihydride and polyhydrides ðSiH2Þn create a

doublet with peaks at 890 cm)1 and 850 cm)1. All of

the bonded hydrogen, however, contribute to the

wagging mode at ~640 cm)1, so the bonded hydrogen

content NHB can be best determined from this peak as

NHB ¼ CB

ZaðxÞx

dx; ð1Þ

where CB is the oscillator strength, x is the optical

frequency of the light, and a is the optical absorption.

The peak in a(x)/x, can be fit to a Gaussian with

amplitude a0 and standard deviation r0, so the total

hydrogen content will be given as

NHB ¼ CBa0r0

ffiffiffiffiffiffi2pp

: ð2Þ

This relationship can then be used to determine

the atomic percentage of hydrogen, assuming an

atomic density of 5.0 · 1022 cm)3, and a CB value of

2 · 1019 cm)3 [19]. The peaks near 2000–2100 cm)1 are

also often used to obtain hydrogen content. The dif-

ferent oscillator strengths for each of the stretching

modes made this calculation difficult for the samples

deposited for this study as many hydrogen-bonding

configurations were observed in the FTIR spectra.

The important role of hydrogen in initiating micro-

crystalline growth is well documented in the literature.

All the growth models rely on the participation of

hydrogen—either through an etching process, a surface

mobility process, or a subsurface chemical process. In

particular, the chemical annealing model relies on a

threshold hydrogen content in the film so that micro-

crystalline sites can nucleate. In Fig. 8, the atomic

hydrogen content of the films grown in this study

is plotted as a function of microcrystalline content.

Despite a fair amount of uncertainty,1 a pattern

indicative of hydrogen-instigated microcrystalline

growth can be seen. As films go from conditions that

generate amorphous films to conditions that result in

microcrystalline films, the hydrogen content increases,

peaks, and then decreases. The decrease occurs as

crystalline sites begin to expand and the chemical

gradient pushes H out of the crystallites, into grain-

boundaries, voids, and towards the surface of the film.

Some of the low-microcrystallinity films show a dis-

proportionally high hydrogen content. This result may

be due to these films being porous or columnar and

having a low density. This theory is supported by a

noticeable Si–O peak near 1000 cm)1, indicating the

deep oxidation of the film, a symptom of low density.

Fig. 6 continued

1 The crystallinity is calculated for films on either SnO2:F, f-SiO2, or CG7059 substrates, whereas the H-content is calculatedfor films on crystalline wafers.


123

Additional information about the films can be

obtained by observation of the stretching mode peaks

around 2000 cm)1. In this area, isolated Si–H bonds

contribute to a peak at 2000 cm)1, clustered Si–H’s

create a peak at 2080 cm)1, and Si–H2 bonds show up

in a peak at 2120 cm)1. To summarize these bonding

configurations, the microstructure parameter R* is

often used. It is defined as the ratio of the integrated

intensity of the peaks above 2000 cm)1 (the peaks

associated with clustered Si–H and Si–H2) to the

integrated intensity of all the peaks in the 1900–

2200 cm)1 range. It gives a good indication of the

quality of amorphous films; films with microstructure

parameters below 0.1 have carrier mobilities that make

them suitable for use in photovoltaic devices.

In Fig. 9, the dependence of the microstructure

parameter R* on microcrystalline content is shown for

the films grown in this study. It can be seen that all of

the films have a fairly high microstructure parameter,

indicating that the hydrogen in the films is clustered in

groups and paired on bonds. Although significant

uncertainty is present, a clear trend can be seen: as the

crystalline content of the films increases, the micro-

structure parameter increases and saturates at ~1. This

Fig. 7 Roughness layerthickness plotted as a functionof microcrystalline fractionXlc for films grown on f-SiO2

Fig. 8 Hydrogen contentversus Xc-Si. The line is aguide to the eye, showing theexpected trend in hydrogen asthe crystalline contentincreases (the line is a guidefor the eye)


123

coincides with the concentration of H at grain bound-

aries and in voids as it is forced out of the crystalline

grains. The fact that this trend can be observed at all is

surprising as all the samples were grown under very

different growth conditions.

Since the microstructure parameter lumps together

the effects of hydrogen singly bonded in clusters and

SiH2, an additional parameter to be addressed is the

proportion of bonded hydrogen existing in clustered

SiH as a fraction of the total integrated intensity of the

peaks around 2000 cm)1. Shown in Fig. 10 is the por-

tion of R* that is contributed by clustered Si–H bonds,

as opposed to Si–H2 bonds and polyhydrides. The

fraction of clustered Si–H goes through a maximum

and then decreases as Xl c increases. This is consistent

with an initial increase in the clustered Si–H as R*

increases due to the bonded hydrogen getting con-

centrated, and then as the hydrogen gets even more

concentrated, an increasing fraction bonds into SiH2

configurations as fewer Si sites are available.

The results of hydrogen content and bonding as

determined by FTIR shows that although the rela-

tionship between substrate bias and microcrystal-

linity is a complicated one, the correlation between

Fig. 10 Dependence ofclustered Si–H fraction (ofR*) on Xlc. Line is guide tothe eye

Fig. 9 Dependence ofmicrostructure factor, R* onXc-Si Line is guide to the eye


123

microcrystallinity and hydrogen content/bonding is

similar to that observed in a-Si:H/lc-Si:H films grown

in other PECVD systems.

3.4 Temperature dependent conductivity and

microcrystallinity

As in most amorphous films, electronic transport

in lc-Si:H films can happen through a variety of

mechanisms, and the dominant mechanism at any

temperature will dictate the behaviour of electrical

conduction. After Kocka et al. [20], we assume the

primary conduction mechanism for these films to be

carriers excited into the conduction and valence bands,

so the resulting measurements were fit to a tempera-

ture-dependent exponential. From this fit, the activa-

tion energy, EA, and the conductivity prefactor, r0, can

be extracted.

rðTÞ ¼ r0eEA=kT : ð3Þ

In Fig. 11, the variation in EA with Xl c is shown for

films grown on f-SiO2. As the microcrystalline fraction

increases, the activation energy approaches the value

characteristic of lc-Si grains due to an upwards shift in

the Fermi level [20]. This is accompanied by a decrease

of several orders of magnitude in r0, as shown in

Fig. 12. In the literature, this has been attributed to a

change in the conductivity path in lc-Si:H due to the

presence of grain boundaries between large grains,

formed when the films become very microcrystalline. It

has been proposed that the increased barriers at the

grain boundaries make the main lateral conduction

path hopping through tails states below the conduction

band, as opposed to extended state conduction. As the

size of the grains increase, the number of barriers

decreases, although their height does not. This explains

the upward trend in r0 whereas EA shows no such

trend.

The results presented in Figs. 11 and 12 clearly show

that the microcrystalline character of the films domi-

nates the conductivity behaviour. The changes in the

lateral conduction mechanism as the film gains crys-

talline character are clearly seen in the conductivity

prefactor and the activation energy of the conductivity.

Once again, controlling the surface potential of the film

during growth has drastically changed the macroscopic

properties of the resulting film.

4 Conclusions

In this work, we have characterized a set of lc-Si:H

films grown by a novel thin-film deposition technology,

the SF-PECVD system. The formation of the micro-

crystalline phase was promoted through the use of an

external electrical bias applied to the growth surface.

The films were subject to PL spectroscopy, AFM, IR

absorption spectroscopy, and conductivity measure-

ments. These results have been correlated with the

microcrystalline content of the films as previously

determined by Raman spectroscopy. In general, it has

been demonstrated that influencing the microcrystal-

line content through the use of substrate surface

Fig. 11 Dependence of EA

on Xlc. As Xlc increases, theactivation energy approachesthat for lc-Si:H. Shown arethe results for films grown onf-SiO2 and the best linear fit isshown


123

electrical bias in the SF-PECVD system has the same

effect on material properties as modifying other

growth parameters in this or other growth sys-

tems—such as the level of hydrogen dilution in

RF-PECVD and HW-CVD systems. We have dem-

onstrated the vital role of substrate bias during lc-Si:H

film growth in SF-PECVD systems in determining all

the resulting film parameters.

Acknowledgements The authors recognize the support of theNatural Science and Engineering Research Council of Canada,ARISE Technologies and Materials and Manufacturing Ontario.

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Fig. 12 Dependence of r0

(log scale) on Xlc. Line is aguide for the eye, showingsimilar results as in [20]


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Physical, electrical, and optical properties of SF-PECVD-grown hydrogenated microcrystalline silicon...

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