P. Roca et al. Tel 33 1 69 33 32 07, Fax : 33 1 69 33 32 07 roca@poly.polytechnique.fr

Microcrystalline Silicon Thin Films Grown by PECVD. Growth Mechanisms and Grain Size Control

Pere ROCA i CABARROCAS, Anna FONTCUBERTA i MORRAL, Billel KALACHE, and Samir KASOUIT

Laboratoire de Physique des Interfaces et des Couches Minces, (UMR 7647 CNRS) Ecole Polytechnique, 91128 Palaiseau Cedex, France.

Keywords : microcrystalline silicon , hydrogen, ion bombardment, in situ ellipsometry

Abstract. Besides plasma reactions, Plasma Enhanced Chemical Vapor Deposition of silicon thin films at low temperatures involves surface and subsurface reactions of the impinging radicals and ions. Dilution of the silane feedstock in hydrogen or high dissociation of pure silane results in a large flux of atomic hydrogen towards the substrate, which can induce a transition from amorphous to microcrystalline silicon growth. In this paper we review previous results based on the layer-by-layer technique, which demonstrate that the ratio of atomic hydrogen with respect to silicon radicals is the main parameter governing the nature of the films and allows for the growth of fully crystallized thin layers on various substrates. These studies highlight the importance of subsurface reactions on microcrystalline silicon formation. We show that the driving force for the formation of stable nuclei is the achievement of a highly porous and hydrogen-rich layer. The plasma is also a source of ions which have varying effects on the film properties, depending on the energy and identity (H+ and SiHx+) of the impinging ions. We discuss the role of ion energy on the different stages of the growth and show that this can be used to control the grain size from a few nanometers up to few tens of nanometers. In particular, we highlight the role of ion energy on the different stages of the growth. Finally, we will present results concerning the use of silicon tetrafluoride as a feedstock, which allow us to achieve polycrystalline silicon thin films ( ~ 100 nm thick) even at a substrate temperature of 200 C.

Introduction

Hydrogenated microcrystalline silicon (mc-Si:H) was first reported by Veprek in 1968 [1]. Due to improvements in the electronic quality of the deposited films, this material has attracted considerable attention recently for use in large area electronics, notably as the absorber layer in p-i-n solar cells [2] or as the channel layer in thin film transistors [3]. However, mc-Si:H films produced by plasma CVD techniques have a complex structure consisting of crystallites of different sizes, an amorphous tissue, and microvoids in variable proportions. In addition, the structure generally evolves as a function of thickness due to the kinetics of crystalline nucleation and competitive grain growth. Therefore one should not consider mc-Si:H as the result of a pure deposition process, but as the result of a growth process driven by physical and chemical reactions. This complex structure has led to the use of different terminologies (nano, micro, or poly silicon) for materials that differ mostly in the grain size - from 3 to 30 nm - and in the crystalline fraction, from ~ 1 - 99 %. The various film structures find application in different fields. At one extreme, nanocrystalline

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films with grain sizes of a few nanometers may exhibit quantum confinement effects when they are embedded in a wide bandgap material such as SiO2 or a-SiN:H. Efficient luminescence can be observed in such nanocrystals provided that their surfaces are perfectly passivated. For this kind of application it would be advantageous to produce films with a high concentration of very small nanocrystalline silicon particles, i.e. to avoid grain growth. We have previously shown that such particles can be synthesized in the plasma and deposited on the substrate [4]. In this paper we will focus on using concurrent ion bombardment during growth as a promising method to produce small crystallites and a high crystalline fraction. Conversely, electronic applications such as thin film transistors and solar cells require grains as large as possible. But, because the grain size generally will be smaller than the film thickness and the channel length in TFT applications, the transport in these materials can be limited by electronic defects at the grain boundaries, which lead to carrier capture or to potential barriers. Fortunately, the plasma deposition process in a hydrogen rich atmosphere at low substrate temperature provides enough hydrogen to passivate the grain boundaries. Detailed studies reveal that transport in these films is mostly limited by the intragrain quality rather than by the grain boundaries [5]. Improving the quality of the grains is a key challenge in the field of large area electronic applications. The goal of this paper is to elucidate the growth mechanisms of mc-Si:H films produced by various RF plasma CVD methods. We will show that in most cases growth involves subsurface reactions and how the resulting film properties result from an equilibrium between hydrogen induced crystallization and ion induced amorphization. We will show how to obtain mc-Si:H films with grain sizes ranging from nano- to micro- to polycrystalline silicon and highlight the effect of ions in the evolution of film growth and on the final material properties

Experiments

The films were produced by the standard, capacitively coupled RF glow discharge technique [6]. For silicon deposition the feedstock was SiH4 or SiF4 and the geometry of the RF electrode was varied in order to tune the energy of the ions. In the layer-by-layer (LBL) approach, the deposition was alternated with a plasma in pure H2. The films were produced at low substrate temperature, in the range of 100 - 250 C. In situ techniques were employed to identify the formation of microcrystalline silicon at different stages of the growth. To determine the film structure, UV-visible spectroscopic ellipsometry data were acquired and modeled using the Bruggeman Effective Medium Approximation (BEMA) and multilayer optical theory [7]. Dark conductivity and contact potential (Kelvin probe) measurements were performed to determine the evolution of the electrical properties of intrinsic and doped mc-Si:H films [8]. In addition, the electron mobility was deduced from time resolved microwave conductivity measurements. Ex situ techniques were used to characterize the structural properties of the films. The hydrogen bonding was characterized via infrared absorption, thermal hydrogen evolution and secondary ion mass spectrometry techniques, while the crystallinity was ascertained by Raman and X-ray scattering techniques. Finally, cross-sectional high resolution transmission electron microscopy was used to characterize the structure of the films as a function of depth and to corroborate the information deduced from spectroscopic ellipsometry measurements [9].

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Results

An important aspect of microcrystalline silicon is the evolution of the structure of the film during growth, first recognized through in situ Raman measurements [10]. This feature, which has been largely studied by in situ ellipsometry [11,12] has also strong effects on the transport properties as deduced from in situ conductivity, Kelvin probe measurements [8], and time resolved microwave conductivity measurements [13]. To simplify plasma chemistry and obtain a better understanding of microcrystalline silicon growth, we have systematically used the layer-by-layer technique, which can be considered a dilution in time with respect to the standard hydrogen dilution method [14]. As discussed below, regardless of the technique or gas precursor, microcrystalline silicon formation proceeds through the same phases. Layer-by-layer deposition. The layer-by-layer technique consists of separating in time the deposition of an a-Si:H film from its exposure to a hydrogen plasma; this combination is referred to as one cycle. Within this approach, mild plasma conditions are used for a-Si:H deposition with typical ion energies below 20 eV [15], and more aggressive ones for hydrogen plasma exposure (in order to achieve sufficient dissociation rates). In previous studies, we showed that for fixed silane and hydrogen plasma conditions the parameter space for microcrystalline silicon deposition can be defined using only two parameters, namely the hydrogen plasma exposure time ,TH, and the silane plasma exposure time ,TSi, which relate to the fluence of atomic hydrogen and SiHx radicals supplied to the film, respectively. For a given value of the a-Si:H deposition time, TSi, the films are amorphous for low values of TH, become crystalline at higher values of TH, and there is no deposition if the hydrogen exposure time is too long and etching dominates. However, this parameter space depends on the plasma conditions as well as on the nature of the substrate, in particular its permeability to hydrogen [16]. It is possible to increase the production of atomic hydrogen and thus decrease TH by increasing the power or the pressure used for the hydrogen plasma [17]. Therefore, it is the deposition rate of a-Si:H which limits the effective deposition rate of the LBL process. While hydrogen is a key element on microcrystalline silicon deposition, its precise role is still a matter of controversy. Among the effects of hydrogen one can consider: selective etching of the a-Si:H tissue, chemical annealing, strained bond breaking and network reconstruction. While selective etching has been suggested as a mechanism for microcrystalline formation, this is not consistent with the LBL experiments where an a-Si:H film is deposited and then exposed to an hydrogen plasma. In order to obtain the nucleation, the formation of a porous hydrogen-rich layer is required. Once this has been formed, the process can evolve towards etching when no SiHx radicals are supplied [18], or towards nucleation and growth of crystallites. Figure 1 shows the imaginary part of the pseudo-dielectric function (e2) deduced from in situ ellipsometry measurements during the LBL growth of mc-Si:H on a thicker film of a-Si:H used as substrate. The data clearly indicate that there is an evolution in the structure of the film during its growth. We see that the spectra recorded at 5 and 10 cycles are characterized by a decrease of the amplitude of owing to an increase of the porosity of the film. The decrease of the amplitude continues up to 15 cycles. At this stage we observe a change, due to the nucleation of crystallites, in the shape of the dielectric function: it becomes flatter and broader in the 3.5 - 4.2 eV energy range. These features are size-broadened traces of the E1 and E2 peaks originating from the crystalline silicon band structure [19]. As the film becomes thicker (e.g. at 25 and 30 cycles in Fig.2), the amplitude of the dielectric function increases and exhibits a more pronounced shoulder around 4.2 eV.

Solid State Phenomena Vol. 93 259

0

5

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15

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25

2 2.5 3 3.5 4 4.5

a-Si:H5 cycles10 cycles15 cycles25 cycles30 cycles

Photon Energy (eV)

Fig. 1. Imaginary part of the pseudo-dielectric function of a mc-Si:H film deposited on a-Si:H by the layer-by-layer method and measured at increasing number of cycles. Note the strong changes in the in the dielectric function, indicating a change in the structure of the film.

The effective dielectric function measured in situ at different stages of the growth, can be decomposed into the volume average contributions of the mc-Si:H structure. Through such analysis, mc-Si:H growth is found to occur through four stages: incubation, crystallites nucleation, grain growth, and steady-state [20]. Incubation. The incubation phase is related to the in-diffusion of hydrogen in the a-Si:H film. We systematically observe that the nucleation of crystallites takes place in a highly porous and hydrogen-rich a-Si:H layer, as illustrated by the HRTEM image in Figure 2a. In-situ studies of the exposure of a-Si:H to a hydrogen plasma have shown that hydrogen can diffuse very rapidly in the a-Si:H network and modify the structure of the film over a thickness of a few tens of nm at 250 C [21,22].

Figure 2. TEM images of microcrystalline silicon films at two stages of their growth using the layer-by-layer technique. The films are deposited on an a-Si:H layer on glass. At seven cycles (left) we can clearly see the formation of a porous layer and the nucleation of some crystallites (darker spots). At 14 cycles the porous layer becomes thinner and denser while the nuclei have grown to larger grains.

Glass

a-Si:H

Nucleation layerPorous layer

50 nm

Glass

a-Si:H

Nucleation layerPorous layer

50 nm

Glass

a-Si:H

Growth

50 nm Glass

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Growth

50 nm50 nm

260 Polycrystalline Semiconductors VII

Nucleation. Once the highly porous (25-30% of voids) and hydrogen-rich layer has been formed the nucleation of crystallites takes place, as shown by the TEM image in Figure 2b. Thus, nucleation of crystallites requires the existence of a poorly linked silicon network that lowers the energy barrier for the transition from a-Si:H to mc-Si:H. Growth and steady-state. Once stable nuclei have been formed, the crystallites grow at the expense of the amorphous phase until the crystalline fraction reaches a steady-state value, which should be viewed as an equilibrium between competing factors. For example, if the hydrogen flux is low compared to that of radicals, nucleation will be reduced and a low crystalline fraction will be achieved at steady state. On the contrary, if the hydrogen flux is too high, a fast crystallization will be obtained but this will result in reduced grain size because of competitive growth of the high number of nucleation sites. The studies of microcrystalline silicon growth by the LBL technique show that that even though the plasma conditions used are these of a-Si:H deposition, a mc-Si:H material can be formed, which clearly demonstrates the role of hydrogen in the nucleation of crystallites. Hydrogen dilution. While LBL is perfectly suited to achieve very thin and highly crystallized layers, it seems less adapted for deposition rates above a few /s. On the other hand, the use of the standard RF at low pressure leads to a severe ion bombardment which, as discussed below, can lead to formation of electronic defects. Thus, to get high deposition rates two approaches have been studied: the use of very high frequencies (VHF) [23] and the use of high pressure [24,25]. In both cases, modeling of the discharge indicates that both parameters lead to similar increase of the electron density and production of reactive species, provided the power coupled to the discharge is kept constant [26].

Figure 3. Effect of pressure and RF power on the deposition rate of microcrystalline silicon films deposited at 200 C from silane-hydrogen gas mixtures. The numbers in brackets indicate the values of the RF power and the partial pressure of silane. The crystalline fraction of the films deduced from spectroscopic ellipsometry measurements was higher than 60 % for all the films.

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Gro

wth

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te (

/s)

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Figure 3 shows the effect of the total pressure and of the RF power on the deposition rate of microcrystalline silicon films. We achieved the increase in deposition rate by throttling the pumping port to increase the silane partial pressure in the silane-hydrogen gas mixture, and by increasing the nominal power at the RF generator. In-situ ellipsometry was used to optimize the process conditions and to characterize the crystallinity of the films. We have been able to increase the deposition rate up to 7 /s while keeping a bulk crystalline fraction above 60 %. However, the increase in deposition rate is accompanied by a thicker incubation layer [17]. Moreover, as in the case of a-Si:H deposition, the increase of the deposition rate is accompanied by an enhanced production of powders in the plasma, which may be the limiting factor for high deposition rates [27]. LBL versus hydrogen dilution. The above results indicate that with respect to the standard hydrogen dilution, the layer-by-layer technique combines a high density plasma required for the production of atomic hydrogen and a soft plasma for a-Si:H deposition. The low energy of heavy SiHx ions in the LBL process clearly reflects on the film properties. This is illustrated in Figure 4 where we compare the structure of thin mc-Si:H films deposited on Corning glass at 250 C. While the imaginary part of the pseudo-dielectric function of the thin films deposited by the LBL and standard hydrogen dilution techniques are quite similar, the optical model used to fit the experimental spectra indicates that despite of being thinner the LBL sample has a higher crystalline fraction and no amorphous interface with the glass substrate.

0

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1.5 2 2.5 3 3.5 4 4.5 5

Photon energy (eV)

Dilution sample

LBL sample

Figure 4. Imaginary part of the pseudo-dielectric function of two mc-Si:H films deposited by the standard hydrogen dilution and the LBL techniques. The multilayer models used to fit the data shown in the right reveal a higher crystalline fraction (Fc) and small amorphous fraction Fa in the case of the LBL sample, as well as the absence of an interface layer with the glass substrate. The complement to 100 % in the interface layer is the void fraction.

Besides crystallinity, the quality of the grains is an important parameter for electronic applications. Figure 5 shows the micro-Raman spectra measured on the same samples as in Figure 4. The spectra have been analyzed assuming three contributions: crystalline silicon, grain boundaries and amorphous phases as shown in Table I. According with the ellipsometry data we found that the amorphous phase is negligible in the LBL sample. Moreover, the position of the crystalline band at 521 cm-1, as in crystalline silicon, and its smaller half width at half maximum (7.5 cm-1) with respect to that of the dilution sample (9.25 cm-1), indicate that the LBL film is less

5.1nm Fc=65% Fv=35% 37.4nm Fc=92% Fa=2% Glass substrate

LBL sampleLBL sampleLBL sampleLBL sample

4.5nm Fc=68% Fa=32% 39.5nm Fc=96% Fa=4% 21.8nm Fc=67% Fa=28% Glass substrate

DDDDilution sampleilution sampleilution sampleilution sample

262 Polycrystalline Semiconductors VII

stressed and has larger grains. The improved crystallinty in the LBL sample, due to the absence of high energy SiHx ions, makes this method an excellent candidate to achieve defect free large grains at small thicknesses as required for TFT applications [28].

200

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400 450 500 550

Ra

man

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sity

(a

.u

.)

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Dilution sample

LBL sample

Figure 5. Raman spectra measured on the samples of Fig. 4. Note that the film deposited by the LBL technique has no amorphous phase, a TO peak at the same position as crystalline silicon indicating a stress free material, and a smaller value of the width, which indicates larger grains compared to the standard hydrogen dilution sample.

SiF4 precursor. Because of the high etching efficiency of fluorine, SiF4 has been widely used as a precursor for the deposition of microcrystalline silicon films from mixtures containing hydrogen [29,30]. Using SiF4, Ar and H2 mixtures we have achieved fully crystallized films without any remaining amorphous phase at the interface, even for layers of a few tens of nanometers [31]. Moreover, the etching effect of fluorine can result in a growth where the incubation phase is absent, i.e. the films are fully crystallized from the beginning. Figure 6 (a) shows the imaginary part of the pseudo-dielectric function of a 130 nm thick film. The optical model used to fit the experimental data using the BEMA approximation indicates that even for such a thin film there is up to 30% large grain polysilicon and this can reach up to 90% for 1 mm thick films [31]. Accordingly, Raman measurements on the same sample shown in Figure 6 (b) are perfectly described by three gaussian peaks at 519.7, 514, and 504 cm-1, without any amorphous component. Moreover, as in the case of the LBL films, we find that the peak position is close to 520 cm-1 indicating a small tensile stress and its width of 5.8 cm-1 is consistent with the achievement of large grains even at such low thickness. These results are also in agreement with the high values of the electron mobility (~ 30 cm2/V.s) measured by time resolved microwave conductivity on microcrystalline films made from SiF4 precursor [13]. Thus, at 200 C using PECVD we obtain materials whose structure is similar to that of polysilicon films produced by solid phase crystallization at 600 C.

cccc----SiSiSiSi GbGbGbGb aaaa----SiSiSiSi

Peak 521 502 485

Width 7.5 34 50

LBL sampleLBL sampleLBL sampleLBL sample

cccc----SiSiSiSi GbGbGbGb aaaa----SiSiSiSi

Peak 525 514 ---

Width 9.25 26 ---

Dilution sampleDilution sampleDilution sampleDilution sample

Table I. Deconvolution of the Raman spectra into three components: crystalline (c-Si), grain boundaries (Gb) and amorphous (a-Si).

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1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0-15

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20a)

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34 1023 264 5 36 50voidsPSi fPSi l

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279 854 8

P h o t o n E n e r g y ( e V )

Figure 6. Spectroscopic ellipsometry (a) and Raman spectroscopy (b) spectra of a thin mc-Si:H film deposited from SiF4 precursor. The inset in Fig. 6a shows the multilayer optical model used to reproduce the experimental data. PSil stands for large grains polysilicon and PSif for fine grain polysilicon. Note the small values of the errors in the evaluation of the film thickness and composition.

Ion bombardment effects. While many studies report that high energy ion bombardment should be avoided when trying to achieve large grains, its role on the growth process and the effect of low energy ions on the material properties have been less studied. To better understand the role of ions on the growth we have co-deposited mc-Si:H films on the anode and the cathode of an asymmetric RF system leading to ion energies in excess of 200 eV on the cathode and below 50 eV on the anode [32]. Figure 7 shows the ellipsometry spectra of two thin films co-deposited on the anode and the cathode from a 2% silane in hydrogen mixture at a substrate temperature of 2500C, an RF power of 20 W, and a total pressure of 900 mTorr, yielding a deposition rate of 0.44 /s and 1.12 /s on the anode and on the cathode respectively. The ellipsometry spectra shown in Figure 7 are treated on the basis of BEMA analysis. The optical model used to fit the spectra - shown on the right of Figure 7 - reveals a film morphology that consists of an interface layer, a bulk layer, and a subsurface layer plus a surface roughness. The interface layer is commonly attributed to a remnant of the incubation phase. The bulk of both samples is fully crystallized though the void fraction is higher for the cathode one. Moreover, the complete crystallization of the bulk on both the anode and the cathode samples means that the hydrogen related crystallization reactions are able to overcome the ion-induced damage. The higher void fraction on the cathode sample is related to a higher hydrogen content due to the implantation of energetic H ions [18]. The subsurface layer exhibits a striking difference for both the composition and the thickness. The lower crystalline fraction and the higher thickness of the subsurface of the cathode sample is consistent with a higher ion-induced amorphization most likely due to heavy SiHx ions, and a higher ion implantation depth. Although sharp interfaces are assumed in the optical model, the real structure certainly will have a gradient of concentration. The differences in the film structure revealed by spectroscopic ellipsometry are further supported by Raman and X-ray diffraction measurements; which indicate a smaller grain size for the cathode film (4.7 nm) relative to the anode film (9.6 nm). Thus, deposition on the cathode allows us to achieve fully crystallized layers with small grain sizes.

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264 Polycrystalline Semiconductors VII

Figure 7. Effect of ion energy on the structure of mc-Si:H films co-deposited on the anode and cathode. Symbols correspond to the experimental data and lines are the result of the model. The multilayer optical model used to fit the experimental data shown on the right indicates that the interface is fully amorphous for the anode sample, while its crystalline fraction reaches 75% for the cathode film. On the contrary the subsurface layer is less crystallized and thicker for the cathode sample.

Discussion: Growth models and grain sizes

Although surface models adapted from a-Si:H deposition studies are still currently used to explain microcrystalline silicon growth [33,34], the evolution of the film properties shown by in situ studies highlights the contribution of subsurface reactions to microcrystalline silicon deposition. Without contradicting such models, H-mediated recrystallization reactions driven by long range diffusion of hydrogen are required to account for the subsurface morphology and the related crystallization. It appears that the growth proceeds through a competition over the subsurface depth between ion-induced amorphization and H-mediated bulk crystallization. Strong support for such a process is given by kinetic studies in which the thickness and composition of the subsurface layer are monitored as a function of the ion energy. Interestingly enough, the ion induced amorphization is found to be an almost instantaneous process while the annealing follows a slower kinetics related to the diffusion process of hydrogen [32]. Based on the above results, we propose the following general schematic picture (Fig.8) of mc-Si:H consisting of three layers plus a surface roughness. As shown in Fig. 8, the interface layer with the substrate may be less crystallized than the bulk, depending on the equilibrium between the deposition rate and the supply of atomic hydrogen. As suggested by the solution of the equation of the hydrogen diffusion [35] and experiments [20], a threshold hydrogen density is required for nucleation to start. Therefore, if the deposition rate is too high or the hydrogen diffusion flux too low the critical concentration for nucleation will not be achieved immediately, but rather after some deposition time thus leaving a partly crystallized interface. Moreover, the requirement of a critical concentration for nucleation can also explain the substrate effect on mc-Si:H growth. Particularly, we observed that it is more difficult to achieve nucleation on a substrate permeable to hydrogen such as a-Si:H compared to glass [16]. Indeed, if hydrogen can diffuse easily into the substrate, to achieve a critical concentration will require an increased hydrogen flux. On the other hand, the high ion energy for films deposited on the cathode can

Interface

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Fc Fa Fv

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Solid State Phenomena Vol. 93 265

increase the critical hydrogen concentration and thus achieve a fully crystallized interface layer as shown in figure 7. Supporting evidence for such process has also been given by spectroscopic ellipsometry measurements in which a flux of fast atomic hydrogen leads to subsurface transformation of a-Si:H to mc-Si:H [36]. The crystalline fraction of the bulk can be ascribed to an equilibrium between the hydrogen-driven crystallization rate, which is related to the diffusion of hydrogen into the film, and the deposition rate. For a fixed hydrogen flux, the increase of the deposition rate will reduce the relative ratio of atomic hydrogen and the crystalline fraction. Finally the subsurface layer clearly is related to the implantation depth of ions. Samples deposited by LBL where the ion energy has been kept low do not exhibit a subsurface layer. On the contrary the subsurface layer is the thickest (40 nm in Fig. 7) for the sample deposited on the cathode by the standard hydrogen dilution. Moreover, a subsurface layer is also present in samples deposited from SiF4-Ar-H2 mixtures (see Fig. 6a), most likely because of the bombardment by Ar+ ions, which have been shown to promote crystallization at low temperatures but leave a less crystallized surface layer [37].

Summary and conclusions

We have presented a review of microcrystalline silicon films prepared by various methods, based on the standard RF glow discharge technique with the aim of highlighting general trends in the growth of this complex material. Because the structure and properties of the material evolve during growth, the use of in situ diagnostic techniques is required for a detailed understanding of the growth mechanisms. Our results show that:

The nucleation of crystallites requires a critical concentration of hydrogen in the a-Si:H layer. Ion bombardment can help to achieve the critical concentration thus providing a way to

produce direct nucleation of crystallites on glass and to avoid the presence of an incubation layer.

Efficient etching with fluorine atoms or a judicious balance between deposition and etching, as in the layer-by-layer technique also achieves a fully crystallized interface.

Ion bombardment appears to be a key parameter to achieve highly crystallized films with small size crystallites.

Figure 8. General layer-structure for microcrystalline silicon thin films. The interface-layer with the glass substrate can be viewed as a remnant of the nucleation process while the subsurface-layer results from an equilibrium between ion-induced amorphization and hydrogen annealing. Note that the crystallization front, which defines the bulk properties, is not at the film surface but below the subsurface-layer.

Glass substrate Bulk

Ion flux H flux

Glass substrate Interface layer

Crystallisation front

Surface roughness

266 Polycrystalline Semiconductors VII

The amorphization of the subsurface layer over a thickness, which depends on the ion energy, strongly support subsurface models in which the crystallization front is located below the ion implantation depth and the annealing of the defects and grain growth is driven by the in-diffusion of hydrogen.

As a conclusion, we have produced microcrystalline silicon films with crystalline fractions ranging from ~ 1 99% and grain sizes ranging from a few nanometers up to a few tens of nanometers. This has been achieved by fine tuning of the plasma parameters, in particular the energy of the ions (deposition on the anode versus the cathode), and the ratio of hydrogen-flux with respect to that of SiHx radicals, easily achieved in the layer-by-layer technique. Moreover, the use of SiF4 feedstock has been shown to result in large grain thin films that favorably compare to polycrystalline silicon produced by low pressure chemical vapor deposition at temperatures above 600 C. While the films with small grain sizes are adapted for luminescent and electroluminescent devices, plasma-deposited microcrystalline silicon films with grain sizes of a few tens of nanometers demonstrates great potential as a new material for macroelectronic devices such as solar cells and thin film transistors.

Acknowledgements

This study was partly supported by the European Community under contract DOIT ENK6-CT-2000-00321,by the Centre National de la Recherche Scientifique, nd the French Agency for environment and energy management (ADEME - Agence pour le Dveloppement et la Matrise de l'Energie ). B. Kalache acknowledges grant from ADEME. The authors acknowledge J.R. Abelson for a critical reading of the manuscript.

References

[1] Z. Iqbal, A.P. Webb and S. Veprek, Appl. Phys. Lett. 36, 163 (1980). [2] H. Keppner, J. Meier, P. Torres, D. Fischer and A. Shah, Appl. Phys. A 69, 169 (1999). [3] I. Chun Cheng and S. Wagner. Thin Solid Films (to be published). [4] P. Roca i Cabarrocas, A. Fontcuberta i Morral, S. Lebib and Y. Poissant. Pure Appl. Chem. 74,

359 (2002). [5] R. Vanderhaghen, S. Kasouit, R. Brenot, V. Chu, J.P. Conde, F. Liu, A. de Martino and P. Roca

i Cabarrocas, J. Non Cryst. Solids 299-302, 365 (2002). [6] P. Roca i Cabarrocas, J.B. Chvrier, J. Huc, A. Lloret, J.Y. Parey and J.P.M. Schmitt. J. of Vac.

Sci. and Technol. A9, 2331 (1991). [7] H. Fujiwara, J. Koh, P.I. Rovira and R.W. Collins. Phys. Rev. B 61, 10832 (2000) [8] S. Hamma and P. Roca i Cabarrocas. Solar Energy Materials and Solar Cells 69, 217 (2001). [9] P. Roca i Cabarrocas, S. Hamma, A. Hadjadj, J. Bertomeu and J. Andreu. Appl. Phys. Lett. 69,

529 (1996). [10] H. Richter and L. Ley. J. de Physique, Colloque C1 43, 247 (1982). [11] R.W. Collins. Appl. Phys. Lett. 48, 843 (1986). [12] M. Fang and B. Drvillon. J. Appl. Phys. 70, 4894 (1991). [13] R. Brenot, R. Vanderhaghen, B. Drvillon and P. Roca i Cabarrocas. Appl. Phys. Lett. 74, 58

(1999). [14] N. Layered, P. Roca i Cabarrocas, B. Drvillon and I. Solomon. Phys. Rev. B 52, 5136 (1995). [15] P. Roca i Cabarrocas, MRS Sym. Proc., Vol. 149, p. 33 (1989). [16] S. Hamma and P. Roca i Cabarrocas. Thin Solid Films 296, 11 (1997).

Solid State Phenomena Vol. 93 267

[17] S. Hamma, D. Colliquet, and P. Roca i Cabarrocas. MRS Symp. Proc. 507, 505 (1998). [18] A. Fontcuberta i Morral and P. Roca i Cabarrocas. J. Non Cryst. Solids 299-302, 196 (2002). [19] H. Touir and P. Roca i Cabarrocas. Phys. Rev B 65, 155330 (2002). [20] S. Hamma and P. Roca i Cabarrocas. J. Appl. Phys. 81, 7282 (1997). [21] M. Otobe and S. Oda. Jpn. J. Appl. Phys. 31, L-1388 (1992). [22] S. Yamasaki, U.K. Das and T. Yasuda. J. Non Cryst. Solids 299-302, 185 (2002). [23] F. Finger, P. Hapke, M. Luysberg, R. Carius and H. Wagner. Appl. Phys. Lett. 65, 2588 (1994). [24] M. Kondo, M. Fukawa, L. Guo and A. Matsuda, J. Non Cryst. Solids 266-269, 84 (2000). [25] T. Roschek, T. Repmann, J. Mller, B. Rech and H. Wagner. J. Vac. Sci. Technol. [26] T. Novikova, B. Kalache, Kh. Hassouni, W. Morscheidt and P. Roca i Cabarrocas. J. Appl.

Physics (to be published). [27] P. Roca i Cabarrocas. J. Non Cryst. Solids 266-269, 31 (2000). [28] P. Roca i Cabarrocas, S. Kumar, V. Tripathi, P. Bulkin, R. Brenot, B. Drvillon, R.

Vanderhaghen and I. French. Solid State Phenomena 80-81, 361 (2001). [29] M. Losurdo, R. Rizzoli, C. Summonte, G. Cicala, P. Capezzuto, and G. Bruno, J. Appl. Phys.

88, 2408 (2000). [30] S. Ishihara, D. He, M. Nakata and I. Shimizu, Jpn. J. Appl. Phys. 32, 1539 (1993). [31] S. Kumar, R. Brenot, B. Kalache, V. Tripathi, R. Vanderhaghen, B. Drvillon and P. Roca i

Cabarrocas. Solid State Phenomena 80-81, 237 (2001). [32] B. Kalache, A.I. Kosarev, R. Vanderhaghen and P. Roca i Cabarrocas. J. Non Cryst. Solids

299-302, 63 (2002). [33] A. Matsuda. Thin Solid Films 337, 1 (1999). [34] J.E. Gerbi and J.R. Abelson, J. Appl. Phys. 89, 1463 (2001). [35] R. S. Crandall, J. Thiesen. Mat. Res. Soc. Symp. Proc. Vol 664, (2001) [36] Y.H. Yang, M. Katiyar, G.F. Feng, N. Maley and J.R. Abelson. Appl. Phys. Lett. 65, 1769

(1994). [37] S. Hamma and P. Roca i Cabarrocas. J. Non Cryst. Solids 227-230, 852 (1998).

268 Polycrystalline Semiconductors VII

Polycrystalline Semiconductors VII 10.4028/www.scientific.net/SSP.93

Microcrystalline Silicon Thin-Films Grown by Plasma Enhanced Chemical Vapour Deposition -Growth Mechanisms and Grain Size Control 10.4028/www.scientific.net/SSP.93.257 DOI References[9] P. Roca i Cabarrocas, S. Hamma, A. Hadjadj, J. Bertomeu and J. Andreu. Appl. Phys. Lett. 69, 29 (1996).doi:10.1063/1.117776 [10] H. Richter and L. Ley. J. de Physique, Colloque C1 43, 247 (1982).doi:10.1051/jphyscol:1982134 [15] P. Roca i Cabarrocas, MRS Sym. Proc., Vol. 149, p. 33 (1989).doi:10.1557/PROC-149-33 [22] S. Yamasaki, U.K. Das and T. Yasuda. J. Non Cryst. Solids 299-302, 185 (2002).doi:10.1016/S0022-3093(01)01160-7 [23] F. Finger, P. Hapke, M. Luysberg, R. Carius and H. Wagner. Appl. Phys. Lett. 65, 2588 (1994).doi:10.1557/PROC-358-745 [24] M. Kondo, M. Fukawa, L. Guo and A. Matsuda, J. Non Cryst. Solids 266-269, 84 (2000).doi:10.1016/S0022-3093(99)00870-4 [28] P. Roca i Cabarrocas, S. Kumar, V. Tripathi, P. Bulkin, R. Brenot, B. Drvillon, R. anderhaghen and I.French. Solid State Phenomena 80-81, 361 (2001).doi:10.4028/www.scientific.net/SSP.80-81.361 [31] S. Kumar, R. Brenot, B. Kalache, V. Tripathi, R. Vanderhaghen, B. Drvillon and P. Roca i abarrocas.Solid State Phenomena 80-81, 237 (2001).doi:10.4028/www.scientific.net/SSP.80-81.71 [37] S. Hamma and P. Roca i Cabarrocas. J. Non Cryst. Solids 227-230, 852 (1998).doi:10.1016/S0022-3093(98)00200-2