Evolution of SiHx hydrides during the phase transition from amorphous tonanocrystalline silicon filmsC. Garozzo, R. A. Puglisi, C. Bongiorno, C. Spinella, S. Mirabella et al. Citation: J. Appl. Phys. 111, 043510 (2012); doi: 10.1063/1.3686136 View online: http://dx.doi.org/10.1063/1.3686136 View Table of Contents: http://jap.aip.org/resource/1/JAPIAU/v111/i4 Published by the American Institute of Physics. Related ArticlesEffect of gold composition on the orientations of oxide nuclei during the early stage oxidation of Cu-Au alloys J. Appl. Phys. 111, 083533 (2012) Surface-induced phase behavior of polymer/nanoparticle blends with attractions J. Chem. Phys. 136, 164904 (2012) In-situ neutron diffraction study of Pb(In1/2Nb1/2)O3-Pb(Mg1/3Nb2/3)O3-PbTiO3 single crystals under uniaxialmechanical stress J. Appl. Phys. 111, 084110 (2012) Magnetic and chemical order-disorder transformations in Co2Fe(Ga1−xSix) and Co2Fe(Al1−ySiy) Heusler alloys J. Appl. Phys. 111, 073909 (2012) Stress-induced VO2 films with M2 monoclinic phase stable at room temperature grown by inductively coupledplasma-assisted reactive sputtering J. Appl. Phys. 111, 073514 (2012) Additional information on J. Appl. Phys.Journal Homepage: http://jap.aip.org/ Journal Information: http://jap.aip.org/about/about_the_journal Top downloads: http://jap.aip.org/features/most_downloaded Information for Authors: http://jap.aip.org/authors

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Evolution of SiHx hydrides during the phase transition from amorphousto nanocrystalline silicon films

C. Garozzo,1 R. A. Puglisi,1,a) C. Bongiorno,1 C. Spinella,1 S. Mirabella,2

R. Reitano,3 S. Di Marco,4 M. Foti,4 and S. Lombardo1

1Consiglio Nazionale delle Ricerche – Istituto per la Microelettronica e Microsistemi, Ottava Strada 5,Zona Industriale, 95121 Catania, Italy2MATIS-IMM-CNR, Via Santa Sofia 64, Catania, Italy3Universita di Catania, Dipartimento di Fisica, Via Santa Sofia 64, Catania, Italy4STMicroelectronics, Str.le Primosole, 50 95121 Catania, Italy

(Received 13 September 2011; accepted 17 January 2012; published online 21 February 2012)

This paper investigates the morphological evolution of hydrogenated amorphous silicon layers

obtained by plasma enhanced chemical vapor deposition at different H dilutions in the regime close to

the formation of the nanocrystalline phase. The role of hydrogen in the transition from the amorphous

to the crystalline phase is investigated by accurate structural and chemical characterization, from the

early stages of nucleation, where the nuclei present size slightly larger than the critical nucleus,

i.e., about 0.8 nm in radius, up to the formation of crystalline grains larger than 30 nm in radius. A

correlation between the structural characteristics of such crystalline phase and the bonding mechanism

of Si with H through multiple hydrides, such as Si-H2 and Si-H3 is found, particularly the tri-hydrides

are found to be directly correlated to the shape and the size of the nanocrystallites present in the

films. The multiple hydrides are found to play a role also in the electrical characteristics of p-i-na-Si:H solar cells whose intrinsic layer is realized in the above H dilution conditions. An explanation

of the experimental data in terms of the different bonding mechanism of H in the Si matrix is provided.VC 2012 American Institute of Physics. [doi:10.1063/1.3686136]

I. INTRODUCTION

During the last three decades and more hydrogenated

amorphous silicon (a-Si:H) has been studied extensively as

base material for thin film solar cells due to the natural abun-

dance of the source material, environmental safety, potential

high performance and capability of low cost production.

However a-Si:H is known to deteriorate with light exposure,

causing a reduction in the power conversion efficiency of the

device. The photoinduced degradation has been attributed to

an intrinsic phenomenon typical of the disordered network

consisting of Si and H, where a nonradiative recombination

of photoexcited electron–hole pairs through a bimolecular

direct process, creates metastable Si dangling-bond defects

(light-induced defects, LIDs) which act as recombination

centers and, thus, decrease the photocarrier lifetime.1–4 The

dangling bonds formation mechanism has been deeply mod-

eled in Ref. 5, by focusing on the hydrogen pairing in the

cases of quenched dangling bonds, corresponding to the clus-

ter phase model,6 or light-induced dangling bonds, corre-

sponding to the H collision model.7 It has been found that

significant improvement in the stability, defined as decrease

of the fill factor after light-soaking, takes place when the

intrinsic layers are prepared from a gas mixture of SiH4

diluted with H2. By increasing the hydrogen dilution, R,

defined as hydrogen to silane flux ratio, H effusion from the

film during the growth takes place8 and the total hydrogen

concentration incorporated in the material is found to

decrease.9 Hydrogen also makes chemical reactions with the

strained Si–Si bonds of the amorphous structure contributing

to the Si–Si bond breaking and atomic rearrangement toward

the crystalline configuration.10 Further increment in the

hydrogen dilution increases the microcrystalline volume

fraction, but the solar cell exhibits a drastic decrease in the

fill factor,9 and the layer photosensitivity can decrease up to

three orders of magnitude.11 The most stable a-Si:H solar

cell has been obtained at a hydrogen dilution just before the

onset of microcrystallinity,12 producing the so called proto-crystalline (pc-Si:H) phase. The pc-Si:H phase is obtained

within a narrow deposition-parameter window, however

within this window the transition between the two phases is

not abrupt, and an intermediate region exhibiting an increas-

ing concentration of small crystallites is present.13 The tran-

sition from amorphous to pc-Si:H phase is also dependent on

the layer thickness. It is known that for substrate temperature

of 200 �C when a thickness of about 250 nm is deposited, if a

dilution of R¼ [H2]=[SiH4]< 15 is used, the layer is amor-

phous, while for R> 15 the layer is completely polycrystal-

line. Much work has been devoted to find a correlation

between the film properties and the gas-phase species in the

plasma,8,14,15 because the dihydride Si–H2 excess bonds are

considered the main responsible for the fill factor light-

induced degradation. Their formation has been attributed to

the formation of higher silane-related-reactive species

(HSRS) in the plasma. HSRS, such as Si4H9, are produced

from silane polymers in the plasma by electron collisions. It

is found that at low electron temperatures, i.e. at high

a)Author to whom correspondence should be addressed. Electronic mail:

rosaria.puglisi@imm.cnr.it.

0021-8979/2012/111(4)/043510/8/$30.00 VC 2012 American Institute of Physics111, 043510-1

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plasma-excitation frequency (>80 MHz), the content of

HSRS and consequently of SiH2 in the plasma decreases.

This paper investigates the morphological evolution of

a-Si:H layers obtained by plasma enhanced chemical vapor

deposition (PECVD) at different H dilutions in the regime

close to the formation of the pc-Si:H phase. The films are

grown at fixed plasma parameters, i.e., RF excitation fre-

quency, power density and substrate temperature, kept con-

stant at 280 �C. This temperature has been chosen to reach a

high deposition rate, necessary to reduce the production

costs, and it represents a range of temperature investigated in

some works present in literature,8 which focuses on the elec-

trical effects of such high rates. The morphological charac-

teristics of the films are followed through advanced

microscopy techniques, from the early stage of nucleation,

where nuclei present size slightly larger than the critical nu-

cleus, i.e., about 0.8 nm in radius, up to the formation of

crystalline Si grains larger than 30 nm in radius. The com-

plex bonding mechanism of H with Si has been deeply char-

acterized, particularly with respect to the formation of

multiple hydrides, such as SiH2 and SiH3. The morphologi-

cal and chemical characteristics have been found to be corre-

lated to the electrical performances, specifically the Fill

Factor, of solar cells fabricated by using the layers of a-Si:H,

with the above H dilutions, as intrinsic layers in a p-i-n struc-

ture. The electrical characteristics have been statistically

measured on 150 solar cells for each H dilution condition.

II. EXPERIMENTAL

Undoped a-Si:H films were prepared by PECVD of a

gas mixture of silane (SiH4) and hydrogen (H2). These films

were deposited on a p type CZ Si substrate 1�4 X cm h100ioriented, with a 50 nm thick silicon dioxide layer deposited

on top. The films were deposited at fixed substrate tempera-

ture equal to 280 �C, by using different SiH4 flow rates, equal

to 100 and 12 sccm. The hydrogen dilution ratio, defined as

R¼FH2=FSiH4, where FSiH4 is SiH4 flow rate and FH2 is H2

flow rate, was varied between 2 and 10, for the layers made

with FSiH4 =100 sccm, and was varied between 10 and 40 for

the other group of samples. Table I summarizes the synthesis

conditions for all the samples, including also the hydrogen

flux percentage defined as H2¼FH2=(FH2þFSiH4) (%).

Moreover, control layers of a-Si 250 nm thick at 40 �Cby magnetron sputtering, and of 50 nm thick SiO2 by

PECVD, have been deposited and used as references for the

RDF, to have a pure a-Si layer with no H, and for the FTIR

analysis, respectively.

To determine the structural characteristics of the de-

posited films we investigated them by high resolution

transmission electron microscopy (TEM) coupled to radial

distribution function (RDF) analysis, by using a JEOL JEM

2010 transmission electron microscope. To avoid sample

crystallization the specimens have been prepared at tempera-

tures lower than 120 �C, and finished by using low energy

ion-milling at 3 keV.

The total H content of the samples has been measured

by elastic recoil detection analysis (ERDA) in glancing angle

geometry. A 2.4 MeV Heþ beam hits the sample surface at

an angle of 75� to the normal, while recoiled atoms (forward

scattered at an angle of 30� from the beam axis) are detected.

A proper Mylar foil placed in front of the detector stops all

the scattered He ions, allowing us to acquire the energy dis-

tribution of the H recoiled atoms. The quantification of H flu-

ences and depth distribution has been carried out by using

two reference samples, with an independently measured H

profile.16 The qualitative analysis of bonded hydrogen of the

samples has been determined by Fourier transform infrared

(FTIR) spectroscopy in the range of frequencies from 400 to

7000 cm�1 and by elastic recoil detection analysis ERDA.

FTIR spectroscopy is used to estimate the content of bonded

hydrogen, through the proportionality constant which corre-

lates the strength of the IR-absorption bands due to the Si-H

vibrations with the hydrogen content. It has been found for

a-Si:H material that the calibration constant of the integrated

strength of the rocking–wagging–rolling vibrations around

640 cm�1 is independent of hydrogen content and of sample

preparation.17 FTIR has been also used to separate the contri-

butions of the several Si-Hx hydrides, such as Si-H, Si-H2,

and Si-H3. Phonon modes for the Si–Hx bonds occur in three

energy bands, a broad peak at 630 cm�1, which is always

present; a group of sharp lines at 800�900 cm�1 whose

shape and intensity depend on deposition conditions; and

modes in the range 2000�2200 cm�1 which depends on the

bonding configuration of the Si with the Hx. At 630 cm�1 all

the three contributions from the Si–H, Si–H2, and Si–H3

bonds are present, while they are well separated in the range

at about 2000 cm�1, respectively at 2000, 2090, and

2140 cm�1.18 The layers of a-Si:H have been used as intrin-

sic layers of p-i-n solar cells, realized by depositing in situdoped pþ and nþ layers, respectively, before and after the

intrinsic layer, on glass substrates covered by transparent

conductive oxide (TCO). The diode formation has been com-

pleted by depositing a second TCO film on the p-i-n struc-

ture, and by defining the device area by lithography.

A number of 150 cells for each H dilution condition has

been electrically characterized in terms of current�voltage

measurements in controlled light conditions, by using a Solar

Simulator with AM1.5 G spectrum, 1000 W=m2 incident

power in steady-state mode, at 25 �C.

III. RESULTS AND DISCUSSION

First the samples were analyzed by Raman spectroscopy

to evaluate the contribution of the amorphous and of the

crystalline phase. The results relative to the samples depos-

ited with FSiH4¼ 100 sccm and with R¼ 2 (H2¼ 66.7%) and

TABLE I. Synthesis conditions of the samples.

R¼FH2=FSiH4 FH2 (sccm) FSiH4 (sccm) H2¼FH2=(FH2þFSiH4) (%)

2 200 100 66.7

10 1000 100 90.0

10 120 12 90.9

20 240 12 95.0

30 360 12 96.7

40 480 12 97.5

043510-2 Garozzo et al. J. Appl. Phys. 111, 043510 (2012)

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R¼ 10 (H2¼ 90.0%) are shown in Figs. 1(a) and 1(b),

respectively. The spectrum relative to the sample deposited

at R¼ 2 has been deconvoluted by a suitable background

model and by using two Gaussian curves centered at 470 and

493 cm�1. The second sample spectrum has been deconvo-

luted by using three Gaussian curves centered at 468, 486,

and 503 cm�1. In literature the first peak, at 468�470 cm�1,

is generally attributed to the position of the TO Raman band

of amorphous silicon, although it is also found that the peak

typical of continuous random network is present at 465 cm�1,

and it shifts to 475 cm�1 when the material presents an

increase in the local order.19 The peak in the range

508�522 cm�1 is attributed to crystalline grains with size

ranging between 2 nm up to microns depending on their size.

Regarding the intermediate peak in the range 480–500 cm�1,

there are different interpretations in literature. It is attributed

either to amorphous20,11 or to intermediate order phases.19 In

the first case such a peak is explained as the breadth of the

amorphous TO phonon band, while in the second case the

490 cm�1 band is explained as the relaxation of the bond-

strain energy correlated to the transition from the disordered

to the relaxed crystalline Si configuration.

To better understand the nature of the thermodynamic

phase of the layers the samples have been analyzed by TEM.

Figure 2 shows the plan view TEM micrographs in dark field

of the same samples as in Fig. 1, i.e., deposited by using

FSiH4¼ 100 sccm with R¼ 2 (H2¼ 66.7%) (a), and with

R¼ 10 (H2¼ 90.0%) (b). In Fig. 2(c) also the TEM image of

the sample deposited with FSiH4¼ 12 sccm and with R¼ 10

(H2¼ 90.9%) is shown. The sample deposited with R¼ 2

was found to be completely amorphous, while both R¼ 10

cases contained a large density of very small nanocrystalline

Si dots approximately of 1.7 nm in average diameter. The

FIG. 1. Raman spectra relative to the samples deposited with FSiH4¼ 100

sccm and with R¼ 2 (a) and R¼ 10 (b).

FIG. 2. Plan view dark field TEM micrographs obtained by selecting the

h111i diffraction ring for the hydrogenated Si layers deposited using

FSiH4¼ 100 sccm with R¼ 2 (H2¼ 66.7%) (a), R¼ 10 (H2¼ 90.0%) (b),

and FSiH4¼ 12 sccm R¼ 10 (H2¼ 90.9%) (c), respectively. The insets in the

Figures represent the diffraction pattern relative to each sample.

043510-3 Garozzo et al. J. Appl. Phys. 111, 043510 (2012)

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insets in the figures report the diffraction pattern relative to

each sample.

The samples obtained at higher hydrogen dilutions

(R¼ 20� 40) were also analyzed by TEM plan view in dark

field conditions and it was found that they were polycrystal-

line with size larger than the previous samples and progres-

sively increasing with H2 dilution. Precisely the samples

with R¼ 20, 30, and 40 presented an average size of about

rd¼ 20 6 1 nm, 21 6 1 nm, and 25 6 1.5 nm, respectively.

The peculiarity of these samples is that the crystalline

regions present a columnar shape when observed in cross

section by TEM, while in the case of both samples deposited

at R¼ 10 no columnar structure has been evidenced, and

they present a spherical shape. The columnar shape of the

samples with R¼ 20� 40 is shown in Figs. 3(a)–3(c) which

reports the TEM micrographs in cross section view, taken in

dark field condition. In the images it is possible to see the Si

substrate (dark gray region), the SiO2 layer (light gray

region) and the deposited Si:H layer on top. This last film

exhibits a columnar morphology with columns shrinking in

size going down toward the substrate.

We can approximate such a shape to a conical geometry.

The height of these cones is approximately equal to

h¼ 200�250 nm in all the three cases, i.e., the crystallized

portion extends for about all the layer thickness. In the

region immediately adjacent at the interface with the SiO2 in

fact the crystallographic phase is not identified. Results pres-

ent in literature showing the same conical structures,21 talk

about a seeding region, made by a mixed phase of amor-

phous and crystal, close to the interface.

The fact that for R � 10 (H2> 90%) we observed the

crystallization of the samples is consistent with the literature

data, which state that there is a critical thickness, dB, above

which the pc-Si:H phase appears.13,22 For thickness of about

200� 300 nm the samples deposited with R< 15

(H2< 92%) should appear in the amorphous phase, while for

higher hydrogen dilutions the samples should be completely

microcrystalline. To explain this apparent discrepancy one

should consider that such work presented in literature

regards samples deposited at substrate temperatures of

200 �C, while the present study refers to samples deposited

at 280 �C. The higher temperature may explain the fact that

the phase transition appears at lower dilutions.

The TEM diffraction patterns from the plan view analy-

sis allowed performing a deep study of the atomic structure

of these nanovolumes through the calculation of the radial

distribution function, RDF. It describes how the atomic den-

sity varies as a function of the distance from one particular

atom. Atomic distributions can be generally investigated by

scattering experiments resulting in either images (e.g., with

electrons) or diffraction patterns (e.g., for x rays, neutrons,

and electrons). However, because of the low scattering cross

section for both neutrons and x rays, the scattered signal is

too small to be used for the study of nanovolumes. In con-

trast, the scattering cross section for electrons is relatively

large and this allows smaller volume of material to be inves-

tigated. An accurate explanation of the calculation of the

RDF analysis can be found in Ref. 23. The calculation of

the RDF started by acquiring the scattered intensity from the

FIG. 3. TEM micrographs in dark field condition obtained by selecting the

h111i diffraction ring for the samples with R¼ 20 (H2¼ 95%) (a), R¼ 30

(H2¼ 96.7%) (b), and R¼ 40 (H2¼ 97.5%) (c).

043510-4 Garozzo et al. J. Appl. Phys. 111, 043510 (2012)

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diffraction pattern,24 I(q) where q is the radially integrated

scattering vector, and is equal to q ¼ 4p sin h=k where 2h is

the scattering angle and k is the wavelength of the scattered

radiation. From the scattered intensity it is possible to calcu-

late the reduced intensity function, u:

uðqÞ ¼ IðqÞ � Nf 2ðqÞNf 2ðqÞ

� �q; (1)

where N is the number of atoms in the investigated volume

and f is the atomic scattering factor. The conversion of I(q)

to u(q) it’s not easy to calculate, because it requires knowl-

edge of the electron atomic scattering factors. However

many researchers23 have noted that useful G(r) can be

obtained by fitting I(q) with a second order polynomial func-

tion, BG(q), as it is frequently done in neutron and x-ray dif-

fraction studies. Peak positions in G(r) are not strongly

influenced by the fit, although this leaves N not univocally

determined. By using this procedure, Eq. (1) becomes:

uðqÞ ¼ IðqÞ � BGðqÞBGðqÞ

� �q: (2)

The radial distribution function, G(r), can be obtained by

Fourier transform of u(q). The calculation was performed in

a fixed integration range of q, between 1 and 8 A�1, in order

to avoid some artifacts present in the high and low frequen-

cies region of the u(q).

Figures 4 (a) and 4(b) show the results of the calculation

of u(q) and of G(r), respectively. In the first figure it is

evident that, by increasing R and H2%, u(q) presents many

peaks, indicating unequivocally that the film evolves

progressively toward c-Si. Moreover the third peak in the

amorphous sample with R¼ 2 is at 5.3 A�1, while for the

sputtered a-Si and for the R¼ 10 samples it is at 5.6 A�1.

This difference in the raw data produces a difference in the

peak positions of the calculated G(r) shown in Fig. 4(b). In

this figure the dashed curves represent the data relative to the

amorphous layers. Particularly the lowest dashed curve indi-

cates the amorphous Si obtained by sputtering, i.e., with no

H dilution, while the dashed curve is relative to the layer de-

posited with R¼ 2. The sequence of curves with continuous

lines indicates the results for the crystalline samples depos-

ited with R � 10. The vertical thin lines at the bottom of the

graph indicate the radius of the 1st to the 7th shell, of Si–Si

bonds in crystalline Si and for the Si–H bonds, as taken from

literature.25

As it is evident in Fig. 4(b), the a-Si:H film (R¼ 2)

presents the first peak at about 1.5 A well evident and sharp,

while in the other samples it is less defined or absent. It is

important to note that in the case of the 250 nm thick control

layer obtained by sputtering, corresponding to the lowest

curve, where it is not expected to have the presence of hydro-

gen, this peak is absent.

This peak can be attributed to the bond between the Si

atoms and their nearest neighbors H atoms, since the Si–H

bonding length reported in literature is equal to 1.5 A. This

sample moreover presents a value of radius at rIshell¼ 2.54

A, much higher than that measured for the amorphous sput-

tered sample (where no H is present). Except for this R¼ 2

case, for the other samples the second and third peaks are

quite close to, or even lower than, the first and second coor-

dination shell in the crystalline Si.25 By increasing R, G(r)

evolves toward that of c-Si, exhibiting peaks that correspond

to the third, fourth and higher crystalline coordination shells.

The appearance of these structural features indicates

unequivocally that the a-Si:H film becomes progressively

more crystalline when R increases. From the above results

we can say that the effect of the H insertion in amorphous

silicon leads to a change in the mean Si–Si bonding length,

as first neighbors, and this effect is visible in the RDF as a

peak typical of the Si–H bond and in the shift of the first

peak toward larger distances.

To understand the correlation between the above results

and the effective content of hydrogen incorporated in the

films, we performed elastic recoil detection analysis (ERDA)

FIG. 4. (a) Reduced intensity, u, as a function of q¼ 4psin(h)=k, and (b) ra-

dial distribution function, G(r), of the films for the several hydrogen dilu-

tions: the samples made with FSiH4¼ 100 sccm, respectively with R¼ 2

(dashed), R¼ 10 (continuous), and the samples with FSiH4¼ 12 sccm (con-

tinuous lines) with R¼ 10, 20, 30, and 40. The data have been vertically

shifted for clarity. The vertical lines in (b) indicate the peak positions of

G(r) for c-Si.

043510-5 Garozzo et al. J. Appl. Phys. 111, 043510 (2012)

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which determines the total hydrogen content, either bonded

or not. The ERDA results are shown in Fig. 5(a).

As it is possible to see in Fig. 5(a) the H content is

almost constant from 66% to about 90%, it presents a peak at

about 91% and then decreases by increasing the H2 flow per-

centage from 95% to 97%.

It should be noted that very small changes in the H2 per-

centage, also 1%, produces evident changes in the H content

in the film. Moreover, the region at about 90% corresponds

to the range where the TEM revealed the formation of the

nanocrystalline spherical dots, while for higher values of H2

percentage the crystallites present a conical shape. The fact

that in the high range of hydrogen dilution the H effectively

incorporated in the film decreases can be correlated to litera-

ture results4,26,27 which present PECVD materials that after

the formation of the pc-Si:H phase the H content decreases

by increasing R. This is due to the fact that H has a role in

the transition from a-Si:H to pc-Si:H and lc-Si, because it

induces a gradual crystallization of the sample, and during

this phase transition a part of H leaves the layer. The mecha-

nism on which is based the H-induced crystallization is

based on the insertion of H atoms into strained Si–Si bonds,

through the formation of intermediate bond-centered

Si–H–Si configurations, as the H atoms diffuse through the

a-Si:H film. After the H moves away from the bond-centered

location, the strained Si–Si bonds either break or relax, then

undergo local structural rearrangements that result in bond

lengths and angles closer to those of c-Si. Under this picture

the H is present mainly on the grain boundaries of the nano-

crystal and not inside them. Figure 5(b) represents a sche-

matic showing this concept. On the left it is shown the

amorphous matrix with the Si-H completely dispersed in

between the Si atoms, while on the right it is reported the

crystalline grain with the H bonds on the boundary. This rep-

resents a simplified schematic of the more complex Si-Hx

configuration which is actually present in the film, as further

shown in the paper.

It should be highlighted that the sample with the highest

content of H does not present any strong peak in the RDF

analysis in correspondence of the Si-H radius [see Fig. 4(b)],

while the sample that exhibits the Si-H peak in the RDF is

the amorphous sample (dashed curve). This may be due to

the fact that, although the amorphous sample has a lower H

content (about half) however the Si in the amorphous struc-

ture is bonded to H as its nearest neighbor (effect to which is

sensitive the RDF analysis), while in the crystallized samples

this “nearest neighbor” effect is lost, because the H is present

only on the boundaries of the nanocrystals.

The schematic of Fig. 5(b) presents the H as a monohy-

dride (Si–H). In the present paper however a deep study on

the bonding configuration of Si–H in terms of multiple

hydrides has been done by FTIR analysis.

Phonon modes for the Si–H bond occur in three energy

bands, a broad peak at 630 cm�1, which is always present; a

group of sharp lines at 800�900 cm�1 whose shape and in-

tensity depend on deposition conditions; and modes in the

range 2000�2200 cm�1 which depends on the bonding con-

figuration of the Si with the Hx. In Table II it is reported the

assignment for the principal vibrational modes of Si-H, Si-

H2, and SiH3 taken from literature.18 As it is possible to see

at 630 cm�1 all the three contributions are present, while

they are well separated in the range at about 2000 cm�1,

respectively at 2000, 2090, and 2140 cm�1.

For this reason the peak at 630 cm�1 has been used to

evaluate the total content of H incorporated in the film, i.e.,

H bonded to Si, and to compare this result to the ERDA data,

which instead provide the H content bonded or not. The

FTIR analysis has also been performed on the reference sam-

ple made by silicon and the same deposited oxide as in the

FIG. 5. (a) Total H dose in the layers plotted as a function of H2 flux per-

centage obtained by ERDA analysis. Also the H2 ratio R is reported for each

result. (b) Schematic representing (left) the Si-H amorphous matrix, and

(right) the nanocrystalline dot with the Si—H bonds on the grain boundary.

This is a simplified schematic of the actual Si–Hx bonding configuration

which is expected to be present in the material.

TABLE II. Si-H vibrational modes.18

Group Frequency (cm�1) Assignment

SiH 2000 Stretch

630 Bend

SiH2 2090 Stretch

890 Scissors

845 Wag

630 Rock

SiH3 2140 Stretch

907 Degenerate deformation

862 Symmetric deformation

630 Rock

043510-6 Garozzo et al. J. Appl. Phys. 111, 043510 (2012)

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layers, in order to properly subtract the signal coming from

the oxide.

The hydrogen content, CH was obtained, as said, by nu-

merical integration of the Si–H rocking–wagging mode at

630 cm�1. The complete procedure is expressed by the fol-

lowing equation:

cH ¼AW

NSi

ð�r=w

að�Þ�

d� � AW

NSi

X�r=w

að�Þ�

d�; (3)

where a(�) is the absorption coefficient of the film at the wave

number �, �r=w stands for the rocking–wagging bands around

630 cm�1, Aw¼ 1.6� 1019 cm�2 is the proportionality con-

stant,17 and NSi¼ 5� 1022 cm�3 is the atomic density of pure

silicon. It has to be noted that depending on the amount of

hydrogen incorporated, the atomic density of the amorphous

Si slightly increases to values of 5.8� 1022 cm�3 for a-Si:H

layers containing hydrogen of around 20 at. %. The results of

the numerical integration are reported in Table III.

For the samples R¼ 2 and both the R¼ 10, since the

total amount of hydrogen is higher than 20%, the value of

NSi¼ 5.8� 1022 cm�3 has been used. As it is possible to see

from Table III the content of bonded H is lower than that

obtained by ERDA, indicating that molecular hydrogen

trapped in internal microvoids or microcavities is present in

the material, in agreement to what previously found in litera-

ture.26 The existence of nanograins in the film provides an

explanation for the loss of H after crystallization: they pro-

vide a connective path to the surface through grain bounda-

ries for the H effusion.28

To deeply analyze the bonding configuration of Si with

H, in terms of mono-, di-, and tri-hydrides, the region at

2000 cm�1 has been analyzed in detail. As seen from Fig.

6(a), by increasing the H2% the whole peak integral

decreases but a tail at larger frequencies appears. The

detailed analysis of such signal, revealed that each sample

presents together three components, at 2000, 2090 and at

about 2140 cm�1, respectively, associated with the stretch

vibrational mode of Si-H, Si-H2, and Si-H3. In the fit proce-

dure the centroids of the three Gaussians have been imposed

to be equal to those of literature, i.e., equal to those reported

in Table II. The results have been normalized with respect to

the total H content, reported in Table III, and successively

respect to each layer thickness as measured by TEM. The

results of this analysis are reported in Fig. 6(b), normalized

respect to the Si-H signal, because it is the largest in all the

cases. The results are shown for the Si-H2 fraction (green

symbols), defined as Si-H2/Si-H, and for the SiH3 fraction

(magenta symbols), defined as SiH3/Si-H. The results are

plotted as a function of H2% (bottom x axis) and of R dilu-

tion (top x axis).

As it is possible to see three regions can be distin-

guished: the first corresponding to the amorphous sample,

where the Si-H2 fraction is prevalent, the second correspond-

ing to the spherical crystallites, for H2¼ 90-91%, where the

Si-H2 fraction is still prevalent, and the third related to

the conical crystals (for H2> 95%) where the contribution of

the Si-H3 fraction is higher than that of the Si-H2.

From the above results it can be concluded that there is

a correlation between the several Si phases and the Si–H

bonding configurations: in the amorphous and in the

crystalline-spherical phases the Si is bonded to the H on the

boundaries of the grain and with a configuration of mono-

and di-hydrides, while in the crystalline-conical case the Si

is bonded to H on the boundaries but as mono- and tri-

hydrides.

TABLE III. Content of H bonded with Si as obtained through FTIR analy-

sis, by numerical integration of the Si–H rocking–wagging mode at 630

cm�1.

R¼ [H2]=[SiH4]

H2¼FH2=(FH2þFSiH4)

(%)

CH

(at. %)

H dose� 1017

(cm�2)

2 66.7 15 2.1

10 90.0 14 1.9

10 90.9 23 3.5

20 95.0 14 1.8

30 96.7 9 1.2

40 97.5 7 0.9

FIG. 6. (Color online) (a) FTIR transmission spectra for silicon hydrogen-

ated layer deposited with the different hydrogen dilution, R, in the

1850�2200 cm�1 range after the background subtraction. (b) Fraction of

Si-H2 (= Si-H2/Si-H) (green symbols) and SiH3 (= SiH3/Si-H) (magenta

symbols) as a function of H2% (bottom x axis) and of R dilution (top x axis).

The lines represent a guide for the eye.

043510-7 Garozzo et al. J. Appl. Phys. 111, 043510 (2012)

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The selective attach of multiple hydrides onto the grain

boundaries of the nanocrystals is close to the literature find-

ings, for which there is a selective attach of multiple

hydrides on the crystallographic orientations (100) and (111)

of Si flat surfaces.29

An interesting correlation has been found between the

fill factor (FF) measured on the p-i-n solar cells and the frac-

tion of Si–H3 bonds measured in the layers. Such a correla-

tion is reported in the graph of Fig. 7, and it is evident that

the fill factor decreases by increasing the trihydrides fraction.

It must be underlined that no correlation is found between

the FF with the Si-H and the SiH2 fraction. This suggests

that the presence of the trihydrides in the intrinsic layer of

the p-i-n is a crucial aspect to monitor and control the final

performances of the solar cells.

IV. CONCLUSIONS

In this paper the structural properties of hydrogenated Si

thin films are investigated in the regime close to the forma-

tion of the protocrystalline phase. The experimental data evi-

denced that the increasing hydrogen flow rate leads to

disorder (a-Si:H) to order (nc-Si) transition, and the average

size of the nanostructures increases. There is a correlation

between the crystallographic phase of the material and the

H2 flow percentage, H2%. Three regions can be identified by

increasing H2%: amorphous, spherical=nanocrystalline

(snXTL), and conical=nanocrystalline (cnXTL). The effec-

tive total H incorporated in the crystalline film increases if

the film passes from amorphous to snXTL phase, while it

decreases when the film goes to the cnXTL phase. In the

amorphous samples the H predominantly incorporated is in

the mono-hydrogen bonding configuration, and this is the

only case in which we are able to detect the nearest neighbor

contribution of the Si–H bond. When the film is in the

snXTL phase the H is present as Si-H and Si-H2, while in the

cnXTL phase the fraction of trihydrides SiH3 is predominant.

The fill factor of the solar cells is found to be strictly corre-

lated to the Si-H3 fraction, and it decreases by increasing this

fraction.

ACKNOWLEDGMENTS

The authors gratefully acknowledge Dr. Nicolo Piluso

and Dr. La Via (CNR-IMM) for the Raman measurements,

and Dr. S. Coffa and Dr. S. Ravesi from STMicroelectronics

for stimulating discussions and support.

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FIG. 7. Relationship between the fill factor of the p-i-n solar cells plotted as

a function of Si-H3 fraction (Si-H3=Si-H). The line represents a guide for the

eye.

043510-8 Garozzo et al. J. Appl. Phys. 111, 043510 (2012)

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