In 1976 Carlson and Wronski announced the first experimental a-Si:H solar cell made at the RCA Laboratory [7]. This single junction p-i-n a-Si:H solar cell deposited on a glass substrate exhibits 2.4% conversion efficiency. It took some years to realize that plasma-deposited amorphous silicon contains a significant percentage of hydrogen atoms bound to the amorphous silicon structure. Hydrogen is essential to increase the electrical quality of the amorphous growth film. Due to the presence of hydrogen, this amorphous silicon film was named hydrogenated amorphous silicon (a-Si:H). The development of a monolithically integrated type of a-Si:H solar cell was of the utmost importance for industrial production [8]. Integration of the monolithic series made it easy to achieve the desired output voltage from a single substrate. In 1980, the integrated type a-Si:H solar cells were commercialized by Sanyo and Fuji Electric, and used in consumer electronics such as

calculators and watches. In the 1980s the initial cell efficiency was increased to around 11-12% [3], which dropped to 7-8% after stabilization. Three decades later, the highest confirmed stable efficiency for a single-junction a-Si:H cell at lab level is 10% obtained in 2010 at the Oerlikon Solar Lab, Neuchatel [9].

Always in the 1980s, industry addressed its interest to off-grid power generation. Modules chosen for this application were single junction p-i-n a-Si:H solar cells. The typical area of the modules ranged from 0.1 to 0.3 m2 and they were designed to deliver a power of around 14 W with up to 5% stabilized efficiency. However, the strong interest in a-Si:H modules cooled down owing to certain limitations that were detected at the time.

The main problem was light-induced degradation known as the Staebler–Wronski effect [10]. This effect is responsible for a typical 15–30% decrease in the initial performance of an a-Si:H solar module. Moreover, inadequate protection and framing of these modules against moisture resulted in the corrosion of contacting electrodes. These problems and the difficulty in further increasing cell (module) efficiency directed researchers to pursue new paths.

In 1983, Matsuda [11] produced both p-and n-type microcrystalline silicon (µc-Si). This new material was used in 1987 by Hattori *et al.* [12] to make the first thin-film silicon solar cell that had microcrystalline-silicon emitters (µc-Si). µc-Si cells attracted general interest mainly because they presented no light-induced degradation. The interest in microcrystalline silicon continued in the 1990s as the market demanded more stable absorbing material, higher output power modules and lower manufacturing costs for thin-film silicon solar cells. New deposition techniques were explored and in 1992 Faraji *et al.* [13] used very-high-frequency PECVD (VHF-PECVD), and in 1997 Rath *et al.* [14] proposed the hot-wire CVD (HW-CVD).

In the 1990s it became quite clear that a-Si:H single junction alone could not compete with low cost crystal silicon and, more recently, with CIS and CIGS-based solar cells. In fact, singlejunction technology is expected to hit its practical efficiency peak (8% to 8.5%) for many producers by 2012, and it will phase out thereafter [1].

To overcome the efficiency limit of amorphous single junction solar cells, industries developed tandem and even triple layer devices containing p-i-n cells stacked one on top of the other. Several a-Si:H based alloys were tested for a better match of the solar spectrum. Hydrogenated amorphous silicon carbide (a-SiC:H) was used as top cell (window layer) [15]. Hydrogenated amorphous silicon germanium (a-SiGe:H) became an attractive low bandgap material for stacked solar cells [16]. Surface textured substrates were also introduced to enhance optical absorption [17].

Some industries transferred these new materials and devices to production lines. For example, a-SiGe:H alloy was used in a tandem solar cell structure by BP Solar [18], Sanyo [19] and Fuji Electric [20], and in a triple junction by United Solar [21]. However, better performances are expected if a higher electrical quality material like µc-Si is used.

In 1994 [22, 23] the "micromorph tandem cell" concept was introduced at the "Institut de Microtechnique" in Neuchatel (Switzerland). It was a new promising thin-film solar cell that used the same deposition technology as amorphous silicon-based devices. The use of the a-Si:H cell combined with microcrystalline silicon (µc-Si:H) leads to the so-called "micromorph" (microcrystalline/amorphous) tandem solar cells.

Because of the difference in bandgap, a-Si layer is suitable for the use of shorter wavelengths and a µc-Si:H layer is suitable for the use of longer wavelengths (Fig. **1**). Therefore a-Si:H/µc-Si:H tandem structure absorbs the sunlight effectively. This technology makes a 30% or more increase possible in output power.

**Figure 1:** Spectral distribution of the AM1.5 solar spectrum incident on the Earth's surface. The spectral ranges which are absorbed by a-Si:H and μc-Si:H, respectively, are indicated.

Moreover, compared to other thin-film technologies currently under development and industrialization (CIGS, CdTe), silicon thin-films have the key advantage of using silicon as raw material which is non-toxic and widely available.

Although the first microcrystalline thin-film was introduced into solar cell fabrication something like 20 years ago, these cells/modules have still not achieved an important increase in conversion efficiency. However, the amorphous/microcrystalline silicon solar cells are still a topic of research and new progress is expected.

### **3. MICROCRYSTALLINE SILICON THIN-FILM DEPOSITION**

Considering that the principal techniques used nowadays for microcrystalline deposition are the same as those used for amorphous deposition, and that they have already been mentioned in this eBook, the easy deposition process will be briefly described below along with the evolution from amorphous silicon to the "micromorphous" concept.

Starting from the initial amorphous silicon deposition and then from the evolved family of semiconductors based on different forms of silicon alloys or crystallinity (nano, micro or poly), all

the various deposition techniques have in common deposition in the form of thin-film at low temperatures (below 600°C and often around 300°C). The general idea of research on µc-Si:H deposition is to understand and control the deposition process in order to achieve the best material quality at the lowest feasible preparation cost. A broad range of vacuum techniques has been employed to work towards this goal.

The three methods commonly used for manufacturing a-Si are electron beam evaporation, rf sputtering and deposition from silane (SiH4) in a radio frequency discharge. Electrical and optical properties of the resulting film depend on the method of deposition.

In these deposition techniques, film growth takes place through sequential sticking of reactive silicon-containing species called growth precursors, which may originate either from a source gas or from a solid target. A common feature of all µc-Si:H deposition techniques is that H2 gas is present to tune the material ranging from a-Si:H to highly crystallized µc-Si:H, as illustrated in Fig. **2**. Crystalline growth starts at nucleation centers near the film/substrate interface. The columnar grains can be several 100 nm wide and several μm long. Atomic hydrogen is believed to play the decisive role in structuring the Si-Si bonds to form a more crystallized matrix [24, 25]. The crystallization process evolves in a transient way, giving rise to a substrate-dependent structural evolution ranging from an initial amorphous incubation zone to the more crystalline thicker parts of the film [26, 27] (Fig. **2**).

**Figure 2:** Schematic diagram showing the prominent microstructural characteristics of µc-Si:H, after Vetterl *et al.* [28]. The left-most region corresponds to highly crystallized material deposited in hydrogen-rich conditions whereas the right-most region corresponds to highly amorphous material deposited in hydrogenpoor conditions (reproduced with permission of Elsevier).

The space between these crystalline grains is filled with amorphous silicon and/or voids. The crystalline volume content of the film depends on deposition conditions, in particular on the silane concentration in the gas mix. The regime around 6% silane concentration is called the "transition region" and yields the best μc-Si:H solar cells [29].

The terms 'micro-crystalline silicon' (μc-Si:H) or, alternatively 'nano-crystalline silicon' (nc-Si:H) are usually adopted to indicate layers that are deposited at low temperatures (below 400°C). The low deposition temperature allows the application of flexible substrates (stainless steel or polymer) that are useful for building integration [29].

Industry was initially drawn to the a-Si:H PV technology as it is a low-T silicon deposition process and the material deposited does not require any post-deposition treatments (such as hydrogenation). This cuts down the thermal budget and saves time.

However, the limited efficiency obtained with amorphous silicon in the 1990s pushed research along new paths which, maintaining these features, could produce solar cells with higher stable efficiency.

Hydrogenated microcrystalline silicon was prepared by Veprek and Marecek for the first time in 1968 by chemical transport method [30]. The fabrication of μc-Si:H using PECVD was introduced in 1980 [31]. The first application of the PECVD technique in microcrystalline deposition was limited to the p type window layers in p-i-n type amorphous silicon solar cells [32]. Hydrogen dilution ratio r = H2/SiH4 > 15 was conventionally used to form μc-Si:H.

The introduction of microcrystalline silicon (μc-Si:H) as i-layer in the 1990s by Wang *et al.* [33] initiated the production of solar cells that possessed higher and more stable efficiency. This material is a heterogeneous blend of a-Si:H, voids and 10 nm c-Si grains. Regardless of the deposition technique used, the resulting material had a very high density of defects.

In 1994 Meier *et al.* prepared solar cells with μc-Si:H i-layer that suffered no light-induced degradation. Later, Klein *et al.* demonstrated that μc-Si:H may show some light-induced degradation with significant a-Si:H volume fractions of 40% or more [34]; however, this degradation is much less pronounced than in a-Si:H and shows improved electronic properties. For this issue, the crystalline volume fraction is a crucial property for materials. Vetterl *et al.*  demonstrated in 2000 that there is an optimum in solar cell performance for material consisting of about 60% c-Si grains, 40% a-Si:H and a minimum amount of voids [28].

The PECVD deposition of silane (SiH4) and silane-containing mixtures ensured better material quality as the result of the incorporation of hydrogen in the silicon matrix that diminishes defects [39]. These features have made PECVD the most widely used preparation technique for thin-film silicon.

**Figure 3:** Diagram of a parallel plate capacitively coupled plasma deposition set up.

The most widespread preparation CVD configuration is parallel plate capacitively coupled plasma deposition using an SiH4 + H2 gas mixture. The parallel plate plasma source consists of two planar electrodes (Fig. **3**). One of the electrodes is the substrate itself on which the silicon film is to be deposited. The other electrode is supplied with an oscillating voltage signal at a typical frequency of 1-100 MHz. The breakdown voltage for plasma ignition strongly depends on the pressure and electrode gap.

For H2 the breakdown voltage shows a minimum of approximately 100 V (about 1 Torr cm). For voltages above the breakdown voltage, ionization reactions ensure the creation of free electrons with a temperature in the range of 104 K and a density in the range of 1010 cm−3 . Due to the inefficient thermal coupling between electrons and the much heavier neutral atoms and molecules, the substrate and background gas temperatures are hardly affected by the hot electrons. The fast electrons create plasma potential compared to grounded walls and substrate over which positively charged ions can gain energy amounting to several tens of eV leading to ion bombardment of the substrate. Moreover, the hot electrons induce chemical reactions that create the reactive precursors involved in film growth. The general growth mechanism and the contribution of the various precursors to µc-Si:H growth is an active area of research.

More details about CVD deposition can be found in dedicated literature but the brief description given here has the sole purpose of underscoring the fact that material properties like crystalline volume fraction are a result of precursor fluxes during deposition. Plasma properties such as electron density and temperature are the result of complex processes like RF power dissipation, inelastic collisions, gas-phase chemistry and surface chemistry. These processes can be empirically influenced by changing process parameters, such as pressure, substrate temperature, RF power, RF frequency, electrode gap, electrode area, gas flow and gas mixture.

Considering the experience developed, it became clear that the guidelines for forming high quality microcrystalline silicon from plasma CVD are very simple as long as a sufficient quantity of hydrogen atoms can be supplied to the substrate during crystallization without causing any damage (most of it is ion damage). Then film fabrication speed can be improved without deteriorating cell characteristics.

Throughput is an important parameter to be considered for mass-production, and it is, therefore, essential to develop a method for high speed film deposition over large areas.

In the case of μc-Si:H, the deposition rate is definitely an issue regarding industrialization as, at present, the thickness of μc-Si:H solar cells is about 5 to 10 times higher than that of a-Si:H solar cells. μc-Si:H solar cells require a thick intrinsic layer (2.0–3.5 μm) to absorb a sufficient amount of sunlight because of its indirect optical transition.

However, thin-film silicon solar cell technology works at relatively low deposition rates, and a high hydrogen dilution condition can limit the growth rate of μc-Si at some A/s. The deposition of a 2 μm thick μc-Si:H film might take 5 h in "standard conditions".

Therefore, high-rate-growth of μc-Si:H film is a crucial matter for low-cost production of thin-film solar cells.

The introduction of very-high-frequency (VHF) PECVD techniques (*i.e.* plasma excitation frequencies well above the standard industrial frequency of 13.56 MHz) has increased growth rates and offers high-quality materials.

The use of plasma excitation frequencies in the Very High Frequencies (VHF) range (70 to 130 MHz) was pioneered by Curtins *et al.* in 1987 [35].

The VHF-PECVD method enabled researchers at the University of Neuchatel to manufacture the first hydrogenated microcrystalline silicon (μc-Si:H) cells at 200°C with reasonable efficiency indexes [36]. These new cells presented maximum efficiency levels of 4.6% and were stable under light soaking conditions. This last feature convinced researchers that it could be possible to produce modules with more stable efficiency levels than the a-Si:H modules.

When a higher excitation frequency is used (typically in the range of 30-300 MHz), electrons in plasma have a much higher thermal velocity than positive ions. This induces a positive potential in the bulk plasma to maintain charge neutrality. The plasma sheath near the substrate leads to the acceleration of positive ions and also to electron confinement within plasma with enhanced dissociation of silane and reduced energy of ions impinging on the growing surface [37]. The use of VHF plasma is a consequence of the improved formation of microcrystalline silicon and higher deposition rate as compared to conventional plasma deposition at 13.56 MHz [38].

In the VHF-PECVD deposition technique the quality of film growth can be controlled by adjusting some deposition parameters.

The thermal budget is another crucial point in crystalline silicon solar cell production. An advantage of PE-CVD is that it allows relatively low deposition temperatures (below 400°C).

During the plasma gas phase, pressure and Silane Concentration (SC):

$$\text{SC} = \text{SiH} \sqrt{(\text{SiH}\_4 + \text{H}\_2)}$$

strongly affect the crystalline and electrical properties of the μc-Si:H film. The structure of the deposited layer changes from amorphous (at high SC-values) to microcrystalline (at low SCvalues). The most efficient solar cells are grown near the μc-Si:H/a-Si:H transition point (SC= 4- 9%).

However, high-rate deposition of device grade μc-Si:H has been generally difficult because of such problems as ion bombardment and powder formation that are often encountered during the high-rate plasma process [40]. In fact, high efficiency μc-Si:H solar cells reported so far can be achieved at limited deposition rates, typically as low as 0.5 nm/s.

An alternative promising deposition technique for the parallel plate plasma deposition of µc-Si:H is the so-called "high pressure depletion" (HPD) regime, [38, 40, 41]. Using this regime, the maximum efficiency of 9.13% has been achieved at an i-layer deposition rate of 2.3 nm/s [42] without important material quality deterioration, and up-scaling to m2 substrate sizes [43] has been demonstrated. Key aspects of the regime are a "small" electrode gap (~1 cm or less), "high"

pressure (~5 Torr or more), "high" power density (~0.5 W/cm2 or more) and "high" H2 dilution (~10% SiH4 in H2 or less). A qualitative view of the µc-Si:H deposition process has been developed to describe the HPD deposition regime. According to this view, dilution in H2 produces sufficient atomic hydrogen for the crystallization process. A high RF power and corresponding high degree of dissociation in the plasma allows high deposition rates to be obtained, but too high RF power gives rise to energetic ion bombardment inducing defects in the deposited material. A high pressure and corresponding short mean free path ensures multiple collisions of ions as they are accelerated towards the substrate, thereby reducing ion energy and preventing defect formation. A small electrode gap is required for plasma operation at high pressure.

Besides the parallel plate plasma deposition, several other deposition methods are being studied.

Growth rates of 3–5 nm/s, which are compatible with industrial applications, can be obtained with the Low Energy Plasma Enhanced Chemical Vapor Deposition [44] techniques see Fig. **4**. This technique is based on a low-voltage, high-current arc discharge which has been so far very successfully applied, in terms of material quality, in the field of Si–Ge hetero-structures grown epitaxially on Si [45].

Preparation of µc-Si:H from a solid silicon target has been reported using reactive sputtering from an Ar + H2 gas mixture [46] or etching and re-deposition (chemical transport) using H2 plasma [30]. Preparation from an SiH4 + H2 gas mixture has been reported using various methods to dissociate the gas, *e.g.* hot wire [47], expanding thermal plasma [48], inductively coupled RF plasma [49], microwave plasma [50-52] and a DC arc [53]. Preparation of µc-Si:H from source gases besides SiH4 + H2 has been reported for *e.g.* SiCl4 + H2 [54], SiH2Cl2 + H2 [55] and SiF4 + H2 [56].

**Figure 4:** Diagram of a LEPECVD deposition system.

Recently, Shinohara [57] reported a variation implemented to the CVS technique, the so-called Localized Plasma Confinement CVD (LPC-CVD) method [58]. The concept of LPC-CVD is to generate a large amount of small plasma by using pyramidal nozzles under very high-pressure conditions, instead of generating uniform large plasma through parallel plane electrodes under conventional conditions. This concept makes it easier to deposit uniform µc-Si:H film with a high deposition rate, because the overlap of the localized plasma can form uniform high-density silane radicals even under high-pressure conditions (>1,000Pa). A maximum deposition rate of 4.1 nm/s was obtained. This feature also helps to increase gas utilization. An a-Si:H/µc-Si:H solar panel with a conversion efficiency of 11.1% (Voc=161.7 V, Isc=1.46 A, F.F.=72.4%, Pmax=171 W), and stabilized efficiency corresponding to 10.0%, was obtained with this technique. The size was 1,100 mm x 1,400 mm and the deposition rate of µc-Si:H was 2.4 nm/s.

However the advantages and disadvantages compared to parallel plate plasma deposition are not clear to date, and more work on these preparation techniques is necessary.

In the laboratory, deposition rates up to 10 Ǻ/s can be obtained for state-of-the-art solar cells with some of these methods, and regular large scale production at such high rates has not been reported as yet, with the exception of Kaneka Corp.

### **4. STRUCTURAL, OPTICAL AND ELECTRICAL PROPERTIES**

The intrinsic layer of a p-i-n (or n-i-p) silicon thin-film solar cell plays two important roles in the functioning process of the device. On the one hand, it must have good light absorption properties in order to generate as many electron-hole pairs as possible for a given incident light intensity and, on the other hand, it must have good electrical transport properties in order to enable the free electron and holes to exit the device before recombining.

Only under certain conditions microcrystalline silicon is suitable for solar cell applications. For instance, it is claimed that favorable conditions for high efficiency pin solar cells are obtained by depositing the μc-Si film at growth conditions near the μc-Si:H/a-Si:H phase transition [62] or close to the transition to amorphous silicon [34].

Other conditions, which must be complied with for high efficiency, are described below:

 the growth rate should be around 1–3 nms-1 to be compatible with industrial applications [40].

In the following paragraph we shall describe the main physical characteristics of undoped and doped μc-Si:H layers in view of their incorporation into solar cell devices.

For such μc-Si:H material, the microstructure varies while the growth of the film proceeds: an amorphous incubation layer has frequently been observed at the bottom (*i.e.* in the initial growth stage) of layers grown with deposition conditions close to the μc-Si:H/a-Si:H transition. Thereafter, the microcrystalline phase follows, consisting of conical conglomerates made of nanocrystals the size of a few tens of nanometers (See Fig. **5**) [59].

The microstructure of μc-Si:H is actually a "mix" between that of amorphous silicon (*i.e.* no periodicity and no long-range order) and that of crystalline silicon (*i.e.* fully regular network of atoms). In fact, three main length scales for disorder can be identified in microcrystalline silicon (see Fig. **5**):

**Figure 5:** Hydrogenated microcrystalline silicon (μc-Si:H).

**Figure 6:** Cross-sectional TEM image of the μc-Si/SiO2/Si region.

Experimental methods allowing the evaluation of i-layer crystallinity are, therefore, of practical importance for the characterization of μc-Si:H. Transmission electron microscopy (TEM) is a very powerful technique that allows for direct observation of the presence of amorphous and microcrystalline phases.

Figs. **6** and **7** show the typical microstructural features of films deposited at low temperature, as obtained from TEM. They present a typical columnar structure that is orthogonal to the substrate surface and a distribution of elongated nanocrystals, with a section of about 5–6 nm, inside the amorphous phase.

**Figure 7:** Cross-sectional HRTEM image of a particular region of μc-Si film. Some stacking faults and some crystalline regions with different crystallographic orientation can be noticed.

Information about the orientation and size of the nanocrystals can be obtained from XRD measurements. Some typical XRD spectra of μc-Si:H samples at different dilutions are reported in Fig. **8**.

In an XRD spectrum of randomly oriented Si powder, a peak intensity ratio of 1:0.55:0.3 among the (111), (220) and (311) peaks is observed.

In this example, the (220) peak intensity is, like the (111) peak intensity, indicative of a preferred (220) orientation of nanocrystallites.

**Figure 8:** X-ray diffraction spectra obtained for a series of μc-Si:H layers deposited on silicon (a) and glass (b) substrates at various silane, d percentage, dilutions. The curves are shifted for clarity. The three main diffraction peaks corresponding to the (111), (220) and (311) silicon crystallographic planes are indicated [60].

The grain size φ along different crystallographic directions can by calculated using the Debye– Scherrer equation:

$$\boldsymbol{\upphi} = (0.9)\,\lambda \langle \text{[K } \cos(2\theta)] \,\text{]}$$

where λ is the X-ray beam wavelength and K is the broadening of the X-ray diffraction line expressed in radiants.

Raman spectroscopy measurements do provide additional information about the microstructural features of the films [61]. Fig. **9** gives an example of the Raman spectrum of sample growth at different dilutions.

By their deconvolution, one could detect the presence of two bands of which the one at about 520 cm-1 is the position of the transverse optical mode in monocrystalline silicon, and the integrated area is associated with the crystalline fraction (IC).

The peak at about 480 cm-1 is associated with the amorphous phase. The integrated intensity of this peak IA is proportional to the amorphous volume fraction in the excited volume.

**Figure 9:** Raman depth profile of a typical CZ set sample grown at d>30%. The crystallinity obtained from each spectrum is reported on the right side. The spectra are vertically shifted for better clarity [60].

The crystalline phase volume fraction χc could be obtained from the Raman spectrum by using the equations:

$$\mathbf{x}\_{\mathbf{c}} = \mathbf{I}\_{\mathbf{C}} (\mathbf{I}\_{\mathbf{C}} + \sigma \mathbf{I}\_{\mathbf{A}})$$

where σ is the ratio between the cross-sections of c-Si and a-Si.

Structural properties have a strong impact on the electrical and optical properties of μc-Si:H films which, again, are closely related to hydrogenated amorphous silicon (a-Si:H), as both materials are deposited by similar techniques. As previously reported, a-Si:H or μc-Si:H material can be

obtained by working on the deposition parameters (such as temperature, silane concentration, power, frequency, pressure…).

μc-Si:H can be considered as something like a (complex) mixture between crystalline silicon (c-Si) and a-Si:H. So it is useful to consider the basic electronic characteristics of a-Si:H before studying μc-Si:H.

As reported in other chapters of this eBook, a-Si:H is a solid-state material made of silicon atoms which are arranged on a lattice that has a certain short range order, but no long range order. Compared to crystalline silicon (c-Si), the average bond angles between neighboring atoms are distorted. Some bonds are even broken and result in so-called "dangling bonds". The presence of hydrogen during the fabrication of amorphous silicon material enables one to passivate a large part of these dangling bonds. These two main "defects" of the lattice of a-Si:H-bond distortion and dangling bonds-give rise to an electronic band structure containing localized states within the socalled "mobility" bandgap (see Fig. **10**). Note that the value of the "optical" bandgap, which is extrapolated from optical measurements, can be slightly different from that of the resistivity or mobility gap.

**Figure 10:** Density of states N(E) for intrinsic a-Si:H. The states are localized (dangling bonds and bandtails) within the mobility bandgap (delimited by EC and EV).

The bond distortion results in band-tails near the valence and conduction bands. In these bandtails, the electrons (or holes) are localized in space and do not participate (directly) in electron transport. The (non-passivated) dangling bonds create deep states near the middle of the bandgap. They can either be positively charged (*i.e.* absence of electron), neutral (*i.e.* one electron) or negatively charged (*i.e.* two electrons). The dangling bonds act as recombination centers for the free electrons and holes. They can also affect electron transport by influencing the total electric charge. In a-Si:H as well as in μc-Si:H, hydrogen has the important role of reducing the number of dangling bonds, or "defects" and, therefore, of making the material "device-grade", *i.e.* suitable for use in optoelectronic devices like solar cells and detectors.

Coming to μc-Si:H films, the co-presence of amorphous and crystalline silicon is mirrored by electrical and optical properties which are a mix of the two materials.

The difference in the absorption range of a-Si:H and μc-Si:H is due to different bandgap energy (Eg) values. For a-Si:H material, Eg ≈ 1.75 eV, whereas for highly microcrystalline silicon Eg ≈ 1.1 eV, similarly to c-Si. Due to its indirect gap, the absorption of μc-Si:H is lower than that of a-Si:H (Fig. **11**).

In addition to the enhanced light absorption, related to its lower bandgap, microcrystalline silicon has an enhanced scattering property due to the natural texture of its surface, which increases effective absorption.

We have already reported that one of the main advantages of μc-Si:H was its stability under light exposure. Works are reported in literature which confirm no light-induced degradation in the transport properties of μc-Si:H layers [63] as well as in the efficiency of μc-Si:H solar cells [22]. In the past, mainly "highly microcrystalline" material, *i.e.* material deposited far from the μc-Si:H/a-Si:H transition, was studied.

However, solar cells exhibiting the highest efficiencies are fabricated near the μc-Si:H/a-Si:H transition. Therefore, light-induced degradation is an important argument for this material and contrasting data can be found in literature. In 1986 Liu [64] reported a light-induced degradation effect for μc-Si:H films containing grain sizes smaller than 120 Å [64]. On the other hand, Fonrodona *et al.* [65] reported the different behavior of nanocrystalline solar cells deposited by HW (hot wire)-CVD at different silane concentrations. In fact, a variation in dangling bond density under illumination was detected for a solar cell deposited at a SC-value of 6.2%. Conversely, they obtained stable solar cells at SC-values of 5 and 5.6%.

**Figure 11:** Typical curves for the optical absorption coefficient of photon energy, for (mono)crystalline silicon (c-Si), as well as for hydrogenated amorphous silicon (a-Si:H) and hydrogenated microcrystalline (mc-Si:H) silicon thin-films [62] (reproduced with permission of Elsevier).

Klein *et al.* [34] have reported up to 10% relative degradation in the efficiency of μc-Si:H HW-CVD solar cells containing a high amorphous volume fraction, after 1000 hours of exposure to light at an intensity of 100 mW/cm2 and a temperature of 50°C.

However, if it seems that the μc-Si:H films are more stable compared to light degradation, on the other hand, μc-Si:H is very sensitive to the incorporation of oxygen, both during deposition and after deposition (post-oxidation). This is why a gas purifier must be used during deposition in order to get "device-grade" microcrystalline silicon material [66]. Fortunately, once incorporated into complete solar cells, μc-Si:H is usually not subjected to post-oxidation. On the other hand, the electrical properties of single intrinsic layers can be dramatically affected by post-oxidation [67], the latter depending on the deposition conditions (especially on the power applied to the plasma). Therefore, the interpretation of transport measurements performed on layers can be a problem.

**Figure 12:** Conductivity as a function of the crystalline volume fraction.

Charge transport in microcrystalline silicon is strongly influenced by its heterogeneous microstructure composed of crystalline grains and amorphous tissue. An even bigger effect on transport is their arrangement in grain aggregates or possibly columns, separated by grain boundaries, causing transport anisotropy and/or depth profile of transport properties.

An example of the impact of crystalline fraction and hydrogen can be seen in (Fig. **12**). Low conductivity values are requested for p–i–n junction fabrication. A decrease in conductivity with increasing crystallinity can be clearly noticed in the figure. This behavior can be explained by the increased presence of grain boundaries inside the film.

Instead, the increase in hydrogen content induces a conductivity increase. This last behavior can be explained assuming a passivation effect of grain boundaries induced by dissolved hydrogen.

Indeed, unlike a-Si:H, where transport models have been successfully developed (see *e.g.* [68]), no simple generally accepted transport model is so far available for mc-Si:H.

Kocka *et al.* [69] recently analyzed electron transport in microcrystalline silicon thin-film and proposed a transport model which considers the columnar structure of this film [70]. A schematic picture of a microcrystalline film is given in Fig. **13**.

**Figure 13:** Schematic drawing of the uc-Si:H film structure and the density of states [69]. (reproduced with permission of Elsevier).

Authors concluded that it is necessary to know the transport path inside the film in order to correctly interpret conductivity. The microstructure of mc-Si:H, namely of the grain boundaries, plays an important role in electron transport. They also formulated a transport model that is based on the idea that there are at least two kinds of c-Si grains, small and LGs. Whenever potential barriers related to "LG boundaries" are formed, the transport properties deteriorate. The detrimental role of LG boundaries on transport stands in contrast with the beneficial optical effect of light trapping *via* scattering on the column heads. This raises the question about what kind of mc-Si:H microstructure is optimal for solar cells, LGs and/or columns. The question is still open.

#### **5. SOLAR CELLS**

### **5.1. Microcrystalline Silicon Solar Cells**

Just like amorphous solar cells, microcrystalline silicon solar cells are usually designed with a p-in structure. The p-i-n structure is necessary as µc-Si:H films have a shorter lifetime than singlecrystal cells. The three layers are a very thin (typically 20 nm) p-type layer, a much thicker undoped intrinsic (i) layer and a very thin n-type layer.

We have already seen that there are, however, some main important differences between μc-Si:H and a-Si:H solar cells; they are summarised below:

 The bandgap of μc-Si:H is about that of monocrystalline silicon (≈1.1 eV at room temperature) whereas that of a-Si:H is significantly higher, *i.e.* around 1.75 eV;

Although μc-Si:H films have these advantages on amorphous silicon material, the higher conversion efficiency so far reached by single junction μc-Si:H solar cells in the laboratory is comparable with that reached by the best single junction a-Si:H cells after degradation (stabilized efficiency) [71].

Moreover, due to the reduced absorption coefficient of μc-Si:H films, a higher thickness is required for optimal light absorption but the deposition of μc-Si:H solar cells takes much longer than that of a-Si:H solar cells and is, therefore, more expensive.

The introduction of the micromorph tandem cell concept [36] has opened new potential. The 1.1 eV (microcrystalline Si) and 1.75 eV (amorphous Si) bandgap combination is very near to the ideal 'theoretical' bandgap combination needed to obtain maximum efficiency under AM1.5 illumination [72, 73].

Single junction μc-Si:H solar cells will be described in the following sections, giving the state of the art for this technology. Then the application of μc-Si:H films in silicon hybrid solar cells will be presented with a tandem structure made of microcrystalline and amorphous silicon thin-films.

*Single junction microcrystalline silicon solar cells*. Thin-film μc-Si:H solar cells can be divided into p-i-n and n-i-p types according to the film deposition order, although the window layer of the solar cell is the p-type layer in both cases. This is done in order to enhance the collection of holes, which are collected at the p-layer, and whose mobility is smaller than that of electrons. A detailed analysis of the transport process in thin-film amorphous silicon solar cells is reported by Deng and Shiff [74]. The asymmetrical drift of electrons and holes explains why amorphous silicon–based pin solar cells are more efficient when illuminated through their p-layers.

Fig. **14** (pin) also reports the two totally different designs for the solar cells involved, corresponding to transparent and opaque substrates connected to the use of the p-i-n or n-i-p configuration.

The n-i-p substrate configuration processing starts with the n-layer, which is normally grown on a metallic back contact. The first n-i-p cells were developed at Kaneka PV Research Division [75] in 1994. These cells had an initial conversion efficiency of 10.1% [75] for a film thickness of 2 µm (surface area 1 cm2 ).

The characteristics of a cell which has a p-i-n structure prepared with the very high frequency glow discharge method were first reported by researchers at Neuchatel University [22]. An initial efficiency of 4.6% and short circuit current densities of up to 21.9 mA/cm2 were obtained.

 (a) (b)

A typical diagram for a p-i-n cell is given in Fig. **14**.

**Figure 14:** Schematic illustration of the structure of p-i-n (a) and n-i-p (b) solar cells.

In the "superstrate" design (Fig. **14(a)**), sunlight enters through the transparent substrate, which is usually glass or a transparent plastic. A conducting layer, typically a "transparent conductive oxide" (TCO) such as SnO2, is deposited on the insulating substrate. The thin-film silicon layers are then deposited on the TCO, starting with a p-type window layer. Finally, a "back" reflector is deposited, which acts as an electrode for the n-type film layer. The "superstrate" configuration is adopted for the p-i-n structure. PV thin-film module manufacturers use soda-lime glass substrates coated with transparent conductive oxide. Soda-lime glass is used for its low cost but this glass has a negative impact on overall cell efficiency due to its absorption in the near infrared. In this spectral region μc-Si:H solar cells still have significant conversion efficiency. A solution could be the introduction in the production line of the more expensive low iron content glass.

In the "substrate" (n-i-p) design (Fig. **14b**), sunlight enters the cell before it reaches the substrate. Starting with the substrate, the cell is manufactured in reverse order, compared to the superstrate design: first a back reflector, then the film layers (starting with an n-type layer), and finally a TCO layer to act as an electrode to the topmost, window layer of the cell.

The main difference between the two configurations is that the underlying layer of a p-i-n cell is the transparent electrode, whereas the underlying layer of a n-i-p cell is the back electrode.

These two designs allow a very wide range of applications for amorphous silicon solar cells. The superstrate design (light enters through the substrate) is particularly suited for building-integrated solar cells in which a glass substrate can be used as an architectural element. The "substrate" design has generally been applied to solar cells using flexible, stainless steel (SS) substrates. The detailed construction of a deposition facility of course depends on whether the substrate is rigid or flexible. Finally, it turns out that the substrate has a profound effect on the properties of the first layers deposited upon it; this effect has led to fairly different cell structures for the "superstrate" and "substrate" designs.

In the thin-film deposition process, the need to keep the device at temperatures not higher than the highest temperature attained during the previous steps is of particular importance throughout the manufacturing process. This means that the "superstrate" configuration allows cheap deposition of SnO2:F by spray pyrolysis (involving temperatures of 600°C to 800°C) at the beginning, so that the 200 to 300°C PECVD process is performed afterwards.

Moreover, the "superstrate" geometry includes an easy encapsulation process and, usually, higher conversion efficiency. In this case, enhanced η values are a common consequence of the relatively simple application of light confinement strategies [76].

Generally, transparent electrodes are made of oxides and there is a risk of these oxides being reduced by the hydrogen atoms used during growth. In the p-i-n structure the oxide window process must be thinner but the reduced film thickness increases contact resistance. The "superstrate" configuration adopted for p-i-n cells has the advantage that cell integration into the modules can be achieved with well consolidated techniques resembling those used for amorphous silicon.

For n-i-p cells, it should be possible to make integrated structures with methods that are equivalent to those used for Cu(In,Ga)Se2-based solar cells [77]. At present, it is hard to determine which is better.

Historically the n-i-p configuration is connected to flexible substrates because it was used on opaque substrates or poorly transparent substrates like steel foils or high temperature polymers. However, the configuration is not limited to this choice. It is, in fact, compatible with any kind of substrates, such as rigid or flexible, transparent or opaque. Nevertheless, flexible substrates are still the main application for n-i-p cells because roll-to-roll processing makes them highly attractive for cutting down production costs as well as energy payback time, particularly when low cost substrates like polyethylene are used. Use of metal substrate requires substrate insulation as, for example, SiO2 films, and this topic is still the focus of research [78, 79].

The schematic band diagram of a p-i-n device is illustrated in Fig. **15**. As with the a-Si:H-based solar cell, a particularly important conclusion that must be drawn is that the absorber layer, the intrinsic one, is indeed the space-charge zone. This means that:

 carrier transport is performed by drift, and not by diffusion; the mobility-lifetime product (µτ) is now the key parameter that determines effective carrier recombination; the average distance covered by a carrier before recombining is given by µτE, where E is the electric field present,

charge zone is not valid. " #

the usual hypothesis (in wafer-based PV) of a negligible recombination in the space-

**Figure 15:** Band diagram of a p-i-n thin-film silicon solar cell.

As already mentioned, by using the AMPS-1D computer program [80] and a set of optimized parameters to describe a-Si:H [81], Schropp demonstrated that it is possible to design electronic levels such as band edges and to gain insight into device physics.

The electric field E(x) within the device causes all electron level energies such as EC and EV to vary in space in the same way. When the device is assembled, these Fermi energies must be equalized to create thermal equilibrium. Electrons are donated from the n-layer to the p-layer, which generates a built-in electric field; while level positions such as EC and EV now vary. The Fermi energy level remains constant.

Carriers are mainly generated in the intrinsic layer, where incoming light generates electron/hole pairs. Both electrons and holes are separated by the internal electric field (drift current), which is built by the doped layers (p-and n-type) on each side of the intrinsic material. Electrons are drifted to the n-type layer, whereas holes, which exhibit lower mobility, are driven towards the p-type one. The geometry presented favors an increased production of electron-hole pairs near the p-i interface, so that hole collection is facilitated by reducing their mean collection distance. It is important to note that carriers generated in the doped layers do not contribute toward the measured current, as the doped material density of defects is too large to allow carrier transport. As a consequence, light absorption in the p-type is to be minimized.

Thus, the electrical, optical and structural properties of the intrinsic layer strongly influence the PV performance of the cell.

Although short circuit current, open circuit voltage and fill factor are all important parameters to be optimized in terms of IV cell characteristics, the open circuit voltage is used here as a guide to overall cell characteristics as it is directly related to the lifetime of the cell. The current is mainly related to light trapping effects and cell film thickness.

In μc-Si:H-based thin-film solar cells, improved performance is achieved [82, 83] when i-layers contain a substantial amount of amorphous tissue but are still in the microcrystalline regime [84]

(a-Si:H/μc-Si:H transition). Bailat *et al.* [59] clearly demonstrated the correlation between the open circuit voltage (Voc) and crystallinity.

Nath *et al.* [85] have reported the dependence of open-circuit voltage (Voc) on crystallinity. Precisely, open-circuit voltage diminishes as the crystalline volume fraction (Fc) increases (Fig. **16**). Authors have identified three contributing factors to the decrease in Voc with increasing Fc, namely an increase in both mid-gap and tail defects as well as a decrease in the mobility gap.

**Figure 16:** The Voc of silicon solar cells with varying Raman crystallinity [86] (reproduced with permission of Elsevier).

The same authors have also [86] reported a deviation to this general rule in fluorinated µc-Si:H solar cells when grain dimensions increase. The large grain fluorinated films grown with high crystalline volume fraction and large grain exhibit an increased performance. They explained this behavior with a reduction in midgap defects and effective density of states at band edges.

The impact of grain dimension was already studied by Werner *et al.* [87]. According to Werner *et al.* the reason why a microcrystalline silicon cell exhibits a typical 10% or more regardless of its small grain size is because of the very low recombination speed at its grain boundaries. This low recombination speed originates from the small barrier height at the grain boundaries, which is guaranteed by (a) the passivation of hydrogen at grain boundary defects, (b) uncharged intrinsic grains (low oxygen density), and (c) the (110) orientation of crystal grains. These hypotheses are in agreement with results obtained by Nath *et al.*

By the close of the 1990s, low-T fabricated μc-Si:H solar cells on glass had reached stable efficiencies of up to 8.5%. However, due to the low deposition rate (<40 nm/min) of the μc-Si:H film and technical difficulties with the development of industrial-scale VHF PECVD systems, μc-Si:H single-junction solar cells do not seem to be commercially viable at present. The same conclusion appears to apply to the 2 μm thick, 10% efficient μc-Si:H STAR solar cell developed at Kaneka using intermediate-T (~550 °C) and standard 13.56-MHz PECVD [88]. Nevertheless, these results showed that a c-Si thickness of merely 2 μm is all that is needed for obtaining a shortcircuit current density Jsc of >24 mA/cm2 .

### **5.2. Contact Layers**

Generally, all thin-film silicon solar cells use a transparent conductive oxide (TCO) layer as front contact, whereas a reflective contact material is needed at the back.

Various transparent conducting oxides have been reported for their solar cell applications. Among the reported TCOs, ZnO [89], ITO [90], and SnO2 [91] are most widely used in the PV area. The PV market requires low-cost solar cells for the widespread use of PV, which can be achieved by introducing low-cost materials and processes for solar cell production. ZnO has been considered as the most promising TCO material in view of cost and properties. Several deposition techniques, such as spray pyrolysis [92], CVD [93] and sputtering [94], have been successfully employed to produce ZnO films. Sputtering has been the preferred method for its production. The window layer in solar cells must be well optimized so that it has low absorption and critical resistivity. The interface properties with underlying layers have to be carefully considered in preparing the window layer.

Very high quality transparent and back reflecting films can reduce the total thickness required for the absorbing layer and, subsequently, the cell cost.

TCO films are strongly affected by deposition parameters (argon ambient pressure and substrate temperature).

For thin-film solar cells based on amorphous (a-Si:H) or microcrystalline silicon (µc-Si:H) with absorber layer(s) in the micrometer range, efficient light trapping is of the utmost importance. Light trapping and effective light incoupling are important aspects for optimized utilization of the solar spectrum. The use of transparent conductive oxides (TCO: ZnO) with a randomly textured surface is a common approach to meet both requirements [95].

As previously reported, two basic cell configurations are used, in practice, for thin-film silicon solar cells: substrates and the p–i–n or 'superstrate' configuration. We shall now see which solution has been proposed to form contact layers for the two configurations.

## p–i–n or 'superstrate' configuration

In the case of a p–i–n deposition sequence, light enters the cell through the glass substrate on which the TCO layer has been deposited as contact layer. We have seen that the p-layer is the first film crossed by incoming light. However we have also seen that the electrical and optical properties of this layer are low. So, this layer should be relatively thin (about 10 nm) and have a low absorption coefficient in order to allow light to reach more efficient layers. The light then crosses the much thicker i-layer (the active layer) and, finally, the relatively thin n-layer (20 nm typical thickness), and reaches (if not yet absorbed) the reflective back contact layer. This last reflective back contact can be a TCO and/or a metal layer. Actual reflectivity is determined by the combination of the TCO interlayer and the metal.

In this 'superstrate' configuration, not only are the electrical and optical properties of the TCO film important, but it is also essential to ensure that it remains chemically stable during plasma

deposition of the other layers. The use of SnO2:F is quite common for a-Si:H solar cells, but must be avoided for single-junction μc-Si:H cells, as it is partly chemically reduced in the H2-rich plasmas that are necessary in this case. Current research also focuses on ZnO, deposited either by sputtering [96] or by low-pressure CVD (LPCVD) [97].

The higher electrical and optical property quality indium-tin-oxide (ITO) film cannot be used in this application as it can be chemically reduced, even without any hydrogen dilution, *i.e.* already by hydrogen formed by silane decomposition during the PECVD process.

A very important aspect for the back contact is its capacity to reflect whatever light reaches this far back into the intrinsic absorber layer. An appropriate combination of a thin TCO layer with a thicker metallic layer (Al or Ag) provides both the desired optical properties (high value of reflectivity) and high electrical conductivity.

## n–i–p or 'substrate' configuration

This configuration has the advantage of using lightweight, unbreakable substrates, such as stainless steel, polyimide or PET. These are generally either opaque, as in the case of stainless steel, or, if they are transparent, they are often easily damaged by ultraviolet light, like most plastics.

Therefore, the substrate/n–i–p configuration is generally used in this case, and lets light enter the device through the semiconductor layer deposited as last layer, *i.e.* still through the p-layer. The top TCO layer is the critical layer, and both ITO and ZnO are used in this case.

Although the deposition sequence differs from the superstrate configuration, the electrical and optical properties required for these films remain stringent.

*Electrical TCO properties* Electrical properties are determined by the sheet resistance of the TCO film, which is related to specific resistivity ρ and film thickness d, *i.e.* Rsheet = ρ/d. Sheet resistance is practically important for integrated serial connection of cells to a solar panel; it gives rise to a series resistance and, hence, to electrical losses.

Obviously, sheet resistance should be minimized; this is achieved by:

Thus, TCO sheet resistance is inherently related to the absorption of near-infrared light within the TCO film. In fact, increasing carrier mobility is the only way to decrease sheet resistance without

**Figure 17:** Example of the impact of Ar pressure and substrate temperature on TCO resistivity. More details about the deposition procedure can be found in the reference [98] (reproduced with permission of Elsevier).

generating a rise in optical absorption of the TCO layer in the near-infrared. The latter may, on the other hand, significantly contribute toward TCO-related optical losses (see below) already at cell level, while sheet resistance causes electrical losses practically only at module level, as a consequence of the integral serial connection of cells.

Considering the mentioned relationship between sheet resistance and TCO absorbance, it follows that efficiency considerations imply a trade-off between these two TCO quantities.

As mentioned, deposition parameters can have a strong effect on the electrical properties of the film. For example, Lee *et al.* [92] have reported the effect of temperature and Ar pressure on film resistivity (See Fig. **17**).

*Optical TCO properties* Apart from the optical absorption that is directly related to sheet resistance, the TCO should have additional optical properties:

Both functions may be achieved by a suitable surface texture of the TCO, with feature sizes comparable to the wavelengths (divided by the refractive index of the corresponding layer).

For shorter wavelengths, the difference in reflectance is due to index grading at the front TCO/player interface. For longer wavelengths penetrating to the back contact, the reduced reflectance stems from both index grading and light trapping.

**Figure 18:** Typical transmission spectrum of TCO film.

A typical transmission spectrum of TCO film is reported in Fig. **18**.

An important aspect should be mentioned here. During the multiple passes of trapped light, light goes through the TCO several times and is reflected by the back contact. Hence, all absorption losses will accumulate, and this makes the need for very low light absorption in contact layers more severe. It is in this field of contact layers and light trapping that most progress is to be expected in cell research, whereas the photoactive absorber layer, the i-layer, has been well optimized.

### **5.3. Micromorph Tandem Silicon Solar Cells**

Although microcrystalline silicon cells formed at low temperatures have a potential for high efficiency, their efficiency in single-cell structures is much lower than that of bulk mono or multicrystalline cells. New concepts have been researched to achieve a substantial improvement in efficiency. In an attempt to achieve this, different groups introduced the concept of two-stacked (hybrid) structures (Fig. **19**) in which multiple cells with different light absorption characteristics are stacked together.

This device structure, as already mentioned, was pioneered by the University of Neuchatel in the 1990s [22].

**Figure 19:** Schematic device structure of the micromorph tandem cell in the 'superstrate' configuration, *i.e.* in the configuration glass/TCO/p-i-n a-Si:H/p-i-n μc-Si:H.

This approach allows better characteristics to be obtained with existing materials and processes. There are different advantages in using a layered structure: it enables a much better use of the solar spectrum; it enables to obtain a higher open-circuit voltage and, it allows, to some extent, suppression of the rate of reduction in cell performance, which is caused by optical degradation phenomena that are observed when using amorphous silicon-based materials (Staebler–Wronski effect). Hence, in principle, higher PV efficiency can be achieved.

Microcrystalline and amorphous silicon are quite good for tandem cells. It can be proven that the bandgap combination of 1.1 eV with 1.75 eV to is very near the ideal combination for tandem solar cells (See Fig. **20**).

*Micromorph Cells Silicon Based Thin Film Solar Cells* **387**

**Figure 20:** Semi-empirical upper limit of efficiency g as a function of the energy gap Eg of the bottom and top cells of a tandem solar cell [99] (reproduced with permission of Elsevier).

Practically speaking, the double layer thickness must be optimized in order to match the current of the two sub cells.

The preliminary structure studied was composed of stacked (tandem) cells in which an amorphous silicon cell was combined with a microcrystalline silicon cell (Fig. **19**).

The low Jsc of stable a-Si:H cells ensures that the thickness of the μc-Si:H cell in the stack does not have to be increased significantly, compared to a stand-alone μc-Si:H cell. Thus, using PECVD, a high PV module efficiency can be achieved with an economic process.

The research team at the University of Neuchatel has coined the term "micromorph" for this tandem cell and reported initial (*i.e.* non-stabilized) efficiencies of up to 13.1% [22].

The "micromorph" concept has also been considered by various organizations, including Kaneka [100]. In 2001 this company started a pilot production of their so-called "hybrid" thin-film PV technology.

The production line is based on low temperature deposition (13.56-MHz PECVD) onto a large (~0.85 m2 ) TCO-coated soda lime glass. The μc-Si:H cell is about 2 μm thick, whereas the a-Si:H top cell is about 5 times thinner.

This company introduced an intermediate thin TCO layer between the amorphous silicon layer of the top cell and the microcrystalline silicon layer of the bottom cell [101].

In this configuration, a transparent intermediate layer reflects part of the light that has not been absorbed while passing through the amorphous silicon layer of the top cell.

As a result, the effective light sensitivity of the top cell is higher than that of a top cell with the same thickness in a conventional hybrid structure. Hence, an improvement can be expected in the short-circuit current Jsc of the overall hybrid cell. The cell with a transparent intermediate layer is more sensitive at shorter wavelengths (top cell).

The presence of the transparent intermediate layer enables to reduce the top cell thickness needed to obtain the same current compared with a conventional hybrid cell. If the thickness of the amorphous silicon film in the top cell can be reduced, then it should also be possible to reduce photo-degradation of the cell.

Quite good initial efficiency rates (aperture area) of over 12% were reported for these large (0.41m2 ) all-silicon modules. The best cell (area 1 cm2 ) made so far has an initial efficiency of 14.5% [102].

Due to the thinness of the a-Si: H cell in the stack, micromorph (or hybrid) PV devices are expected to decay very weakly in the field. The best stable efficiency confirmed as yet is 11.7% for a sub-module [103] produced by Kaneka in 2004 [104], and 11.9 for a cell manufactured by the Neuchatel group in collaboration with industrial partner Oerlikon [105].

These technologies promise high stable average efficiency of about 10% for large-area massproduced modules using large-area low-T processing equipment from the LCD industry.

However, they require a textured TCO layer and a high capital cost of the deposition tool for the ~2 μm thick μc-Si:H cell, which can be a limitation for high throughput production lines.

### **CONCLUSIONS**

The introduction of microcrystalline silicon in thin-film silicon solar cells is relatively recent. The interest in microcrystalline silicon issued from the need to overcome some intrinsic limitations of amorphous solar cells. However the results so far obtained are still unsatisfactory and, in fact, incorporating μc-Si:H is, thus, still a topic of research.

Several deposition techniques have been tested, and the PE-CVD ones have achieved the best results so far. PE-CVD equipment is already used for a-Si:H deposition with a mixture of silane and hydrogen, and increasing the hydrogen to silane ratio in the feed gas easily obtains microcrystalline growth with reduced variation of equipment.

This technique has many advantages, such as low temperature deposition, use of non-toxic materials and deposition on large substrates. All these features offer the potential for future cost reduction of photovoltaic modules based on this material. Hydrogenated microcrystalline silicon is a complex material composed of a mixture of crystalline and amorphous silicon that can be more easily doped than a-Si:H. On the other hand, it is also more sensitive to contaminants than a-Si:H (for example, oxygen). This material exhibits a wide range of microstructures that depend both on deposition conditions and on the substrate material.

It is quite clear that the best solar cells are obtained when the deposition conditions (hydrogen dilution, temperature,) are tuned in order to approach the microcrystalline/amorphous transition.

The strong interest in μc-Si:H also comes from the quite different optical properties of μc-Si:H compared to a-Si:H. Although μc-Si:H has an indirect optical bandgap at 1.1 eV and has, therefore, a relatively low optical absorption, this optical gap matches the higher bandgap of a-Si:H (about 1.75 eV) for a tandem cell quite well.

The electron transport properties of μc-Si:H resemble those typical of a-Si:H but with increased conductivity. The carrier transport mechanism can be explained considering the amorphous/disordered material present in μc-Si:H layers.

Although the electronic properties are quite better, lifetime is reduced and a p-i-n type or n-i-p type device structure is used also for microcrystalline silicon-based devices. In such devices, carrier collection is enhanced by drift, *i.e.* by the internal electric field that prevails in the i-layer. Currently, the conversion efficiency of the best single junction μc-Si:H solar cell is around 10%.

Microcrystalline silicon was initially introduced in single junction μc-Si:H solar cells due to the reduced electrical performance degradation associated with the Staebler–Wronski effect. However, the conversion efficiency achieved so far by the best single junction μc-Si:H solar cell is around 10%.

One expects a further increase in the efficiency of μc-Si:H single junction solar cells mostly by implementation of improved light trapping schemes and by a corresponding increase in short circuit current density. Improved light trapping will be based on the development of new and better suited transparent conductive oxides, structured substrates (periodic gratings or rough random substrates).

More improvements are expected using microcrystalline silicon in tandem cells consisting of a stack with a top a-Si:H cell and a bottom μc-Si:H cell. This combination of materials can be shown theoretically to possess bandgaps that are very near to the best combination of bandgaps for optimum conversion of the AM1.5 spectrum. For reaching maximum efficiency, an intermediate reflector is used between the bottom and the top cell. Such tandem devices generally suffer less from light-induced degradation in the amorphous top cell. In general, the microcrystalline bottom cell inside the tandem is stable (as it is only exposed to the long wavelength part of the solar spectrum). In order to gain both in conversion efficiency and in device stability, one should reduce the thickness of the amorphous top cell and further optimize light trapping. This may hopefully lead to solar cells with stabilized efficiencies of 13–14%. The industrial production of tandem modules has currently been started by a Japanese module manufacturer.

Several other firms (mainly Japanese) are carrying out research on this subject and a few of them have even developed prototype modules.

Thin-film silicon offers interesting applications in building-integrated photovoltaics through its esthetical appearance and the versatility provided in the choice of a substrate material. On the other hand, the 'micromorph' tandem concept allows for an increase in module efficiency towards a stabilized value above 10%. Because of these factors, one may reckon that the industrial development of thin-film silicon tandem modules based on a combination of amorphous and

microcrystalline silicon will be intensified during the next few years, leading to a clear reduction in module costs, probably down to values below 1\$/Wp.

### **ACKNOWLEDGEMENT**

Declared none.

### **CONFLICT OF INTEREST**

The author(s) confirm that this chapter content has no conflict of interest.

#### **REFERENCES**

© 2013 The Author(s). Published by Bentham Science Publisher. This is an open access chapter published under CC BY 4.0 https://creativecommons.org/licenses/by/4.0/legalcode

**CHAPTER 11**

## **Light Trapping in Thin Silicon Solar Cells**

**Mario Tucci1,\*, Luca Serenelli1 , Simona de Iuliis1 , Domenico Caputo2 and Giampiero de Cesare2**

*1 ENEA, Research Center Casaccia, Via Anguillarese 301, 00123 Roma Italy and 2 DIIET, University "Sapienza", Via Eudossiana 18 00184 Rome, Italy*

**Abstract:** Light trapping plays a key role in solar cell to enhance light confinement and then light absorption within the cell. The reduction of the optical losses reflects in photovoltaic solar cell efficiency enhancement. To this aim different and combined strategies can be adopted. Firstly a reflectance reduction of sunlight impinging on the cell is mandatory. Then a cell surface texturing can be very helpful to reduce the reflection by increasing the chances of reflected light bouncing back onto the surface rather than surrounding air. To this purpose different approach can be followed such as substrate and/or single layer texturing. Moreover particular care should be paid to the rear side of the cell where the introduction of a reflecting mirror can produce an optical path length enhancement over a wide wavelength spectrum. Rightly combining the front side texturing and the rear side mirroring it is possible to enhance this pathlength up to 50 times the device thickness indicating that light bounces back and forth within the cell many times performing a light confinement. To this purpose Bragg reflector formed using thin film technology is one of the most promising approach to the back side reflector. Also multi-junction cell, such as tandem micromorph, can receive benefit to the light management by the introduction of inner reflector within the two cells to enhance the spectral separation between the two cells and reduce the front cell thickness where the metastability problem still remain an open request.

In this work we overview all these concepts as applied to thin film cell and to silicon based cell that actually are speedily moving toward thinner substrate making the light confinement a key point to cell efficiency enhancement and PV cost reduction.

**Keywords:** Amorphous silicon, antireflection coating, heterojunction, lifetime, numerical model, optical properties, reflector, scattering, solar cell, texturing, thin film, transparent conductive oxide, plasma process, light trapping, thin wafer, quantum efficiency, bragg reflector, haze, reactive ion etching, laser treatment, silicon.

### **1. INTRODUCTION**

To enhance the solar cell efficiency in recent years, researchers have experimented with approaches to capture the sunlight photon impinging on the device and forcing it to collaborate instead of being reflected away. This is done by changing the basic structure of solar cells, so that photons become trapped when they arrive in contact with the cell, thus providing a greater chance that they will be absorbed and converted into electrons. To this aim a light-trapping scheme introduction is mandatory. Light trapping allows cell to absorb sunlight using an active material

**\*Address correspondence Mario Tucci:** ENEA, Research Centre Casaccia, *Via* Anguillarese 301, 00123 Rome Italy; Fax: 0044 1740 644146; E-mail<EMAIL_ADDRESS>
layer thinner than the material's intrinsic absorption length. This can reduce the amount of materials used in PV cells, cutting cell cost that are mainly based on less abundant or expensive purified materials. It is well known that the largest part in cost production of industrial silicon solar cells is actually referred to the base silicon substrates. Hence reducing the amount of involved silicon, by means of thinner wafer, is a key point in limiting the overall PV cost, making the price of energy from solar cells competitive with the conventional production systems [1]. Moreover, light trapping can improve cell efficiency, since thinner cells provide better collection of photo-generated charge carriers [2]. Historically the theory of light trapping was developed for conventional cells in which the optical path length is several times the actual device thickness, where the optical path length of a device refers to the distance that an unabsorbed photon may travel within the device before it escapes out of the device [2, 3]. For example, a solar cell with no light trapping features may have an optical path length of one device thickness, while a solar cell with good light trapping may have an optical path length of 50 times the device thickness, indicating that light bounces back and forth within the cell many times. Basically conventional light trapping exploits the effect of total internal reflection between the semiconductor material, such as silicon, and the surrounding medium commonly assumed to be air. Simply by texturing the semiconductor surface it is possible to couple light obliquely into the medium, thus giving a longer optical path length than the physical device thickness and also it is possible to randomize the light propagation directions inside the material. Achieving a total internal reflection results in a much longer propagation distance inside the material and hence a considerable absorption enhancement. Indeed the light can be trapped inside the cell and make multiple passes through the cell, thus allowing even a thin solar cell to maintain a high optical path length. The standard theory for this light-trapping scheme was developed by Yablonovitch [3] and applied to silicon based solar cell several years later [4]. To further enhance the absorption of light penetrated into the device a rear side reflector is needed to avoid light escape from the rear side of the cell. This allows the light to bounce back into the cell for possible absorption. This allows an optical path length of approximately 50 times the physical devices thickness and thus is an effective light trapping scheme. For thin film cells with thicknesses comparable or even smaller than wavelength scale, the ray-optics representation and some of the basic assumptions adopted in the conventional theory are no longer applicable. Recently in literature several numerical [5-8] and experimental [9-17] works reported considerable improvement beyond the Yablonovitch limit.

In this chapter a description of light management at the front and rear side of the solar cell is reported. In this description not only thin film cell are considered, but also silicon base cells that are going beyond thinner silicon wafer to reduce PV cost. Indeed in this case several approaches can be helpfully transferred from thin film technologies to wafer based technology. In this direction the best silicon based cell candidate is the amorphous/crystalline heterostructure. To exploit this argument a particular section is dedicated to a dielectric reflector and its related manufacturing technology helpful to improve thin solar cell performances. All the technologies that can be proposed to the light confinement must fit with industrial manufacturing facilities otherwise can only represent academic speculations not useful to reduce the cost per watt produced by the sunlight photovoltaic conversion. Therefore the use of sophisticated technology such as photo-lithographic steps, not easy to transfer to very large area device, used to produce particular features on front or rear side of the cell to enhance the light confinement are not extensively exploited in this overview.

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **397**

In this chapter a detailed description of cell design is reported concerning the use of antireflection coating and texturing of both thin films and silicon wafer based device in order to manage at the best the light impinging on a solar cell. After an overview of light scattering promoted by texturing a detailed description is proposed on rear side reflector in which a Bragg scheme is detailed and experimented.

### **2. ANTIREFLECTION COATING**

Optical losses are responsible for the short circuit lowering in solar cells. They are mainly caused by the parts of incident light escaping from the cell and the part reflected from the cells surface or metal front contacts. To maximise the percentage of light to be converted in electron-hole pairs it is important to maximise the light penetrating into the cell. This can be done modifying the surfaces morphology producing the so-called texturing and applying to the front surface an Anti Reflection Coating (ARC) [18]. The idea behind anti-reflection coatings is that the creation of a double interface by means of a thin film gives you two reflected waves.

**Figure 1:** Schematic of Anti Reflection Coating principle of operation for a single layer ARC.

If these waves are out of phase, they partially or totally cancel. In practice a thin layer of dielectric material having a precise refractive index is deposited with a specific thickness on a substrate, so that interference effects in the coating cause the wave reflected from the anti-reflection coating top surface to be out of phase with the wave reflected from the semiconductor surfaces. In particular when the light reflected from the second interface arrives back at the first interface 180° out of phase with that reflected from the first interface, there is no net light coming out, at least for a chosen wavelength. The schematic situation is sketched in Fig. **1**, where *n0*, *n1* and *n2* are surrounding ambient, dielectric film and semiconductor refractive index respectively. The expression for the fraction of the energy in a normally incident beam of light reflected from the surface of a material covered by a transparent layer of thickness *d1* is [19, 20]:

$$R = \frac{r\_1^2 + r\_2^2 + 2r\_1r\_2\cos 2\theta}{1 + r\_1^2r\_2^2 + r\_1r\_2\cos 2\theta} \tag{1}$$

with:

$$r\_1 = \frac{n\_0 - n\_1}{n\_0 + n\_1} \ r\_2 = \frac{n\_1 - n\_2}{n\_1 + n\_2} \ \ \ \theta = \frac{2\pi n\_1 d\_1}{\lambda} \tag{2}$$

In this case the reflectance has its minimum value:

$$R\_{\rm min} = \left(\frac{n\_1^2 - n\_0 n\_2}{n\_1^2 + n\_0 n\_2}\right)^2 \tag{3}$$

which can be driven to zero when 1 2 0 2 *n nn* , for a specific wavelength *λ0*, due to the dependence of the refraction index on wavelength.

**Figure 2:** Reflectance profiles simulations of Silicon Nitride on flat silicon as a function of silicon thin film thickness. The Black continues line is the AM 1.5 G spectrum, expressed in arbitrary units.

However the sun spectrum is of course made by multiple wavelength photons, so the choice of material, refractive index and thickness of the dielectric film needs to be done considering the best reflection reduction around the wavelength corresponding to the maximum of the sun spectrum intensity. For photovoltaic applications, the refractive index and thickness are chosen in order to minimise reflection for a wavelength of 600 nm [21]. At the present one of the most widely used material is Silicon Nitride deposited *via* Plasma Enhanced Chemical Vapour Deposition (PECVD) which refractive index can be varied from 1.5 and 2.3 circa, evaluated at 633 nm [22]. In Fig. **2** we show simulations made with the aid of NKDGEN software [23] of how the reflectance profile (in term of % of incident light) varies with the thickness of a silicon nitride film (as deposited on a flat

silicon surface) having a refractive index of 2 at 633 nm. It is appreciable the best profile giving the lower reflectance around 600 nm is around 60-70 nm thickness.

Such a kind of behaviour can be enhanced with the introduction of a multi-layer as ARC (see Fig. **3**), varying the refractive indexes and thicknesses to obtain a low reflectance profile over a wider range of wavelengths. The equations for multiple antireflection coatings are more complicated than that for a single layer [20]. Any film combination composed of two dielectric layers on a semi-infinite substrate has a reflectance given by:

$$R = \frac{r\_1^2 + r\_2^2 + r\_3^2 + r\_1^2 r\_2^2 r\_3^2 + 2r\_1 r\_2 (1 + r\_3^2) \cos 2\varphi + 2r\_2 r\_3 (1 + r\_1^2) \cos 2\varphi + 2r\_1 r\_3 \cos 2(\varphi + \varphi) + 2r\_1 r\_2^2 r\_3 \cos 2(\varphi - \varphi)}{1 + r\_1^2 r\_2^2 + r\_1^2 r\_3^2 + r\_1^2 r\_3^2 + r\_2^2 r\_3^2 + 2r\_1 r\_2 (1 + r\_3^2) \cos 2\varphi + 2r\_2 r\_3 (1 + r\_1^2) \cos 2(\varphi + \varphi) + 2r\_1 r\_2^2 r\_3 \cos 2(\varphi - \varphi)} \tag{4}$$

with:

$$r\_1 r\_1 = \frac{n - n\_1}{n + n\_1}, \ r\_2 = \frac{n\_1 - n\_2}{n\_1 + n\_2}, \ r\_1 = \frac{n\_2 - n\_3}{n\_2 + n\_3}, \ \ \wp = \frac{2\pi n\_1 d\_1}{\lambda} \quad \ \ \wp' = \frac{2\pi n\_2 d\_2}{\lambda} \tag{5}$$

where *n* is the refractive index, *d* the film thickness and the wavelength. This expression is quite cumbersome and can be solved only by approximate numerical methods, unless the two layers are restricted to be of equal optical thickness at 550nm. In this case the above equation can be simplified to:

$$R = \frac{r\_1^2 + r\_2^2 + r\_3^2 + r\_1^2 r\_2^2 r\_3^2 + 2r\_2(r\_1 + r\_3)(1 + r\_1 r\_3)\cos 2\varphi + 2r\_1 r\_3 \cos 4\varphi}{1 + r\_1^2 r\_2^2 + r\_1^2 r\_3^2 + r\_2^2 r\_3^2 + 2r\_2(r\_1 + r\_3)(1 + r\_1 r\_3)\cos 2\varphi + 2r\_1 r\_3 \cos 4\varphi} \tag{6}$$

For the equal optical thickness of quarter-wavelength coatings (*i.e.* 11 2 2 0 *nd nd* / 4 ) and in terms of the indices media *n0* of ambient, *n1* and *n2* for two coatings and *n3* for substrate, R at λ0 becomes

$$R = \left(\frac{n\_1^2 n\_3 - n\_2^2 n\_0}{n\_1^2 n\_3 + n\_2^2 n\_0}\right)^2 \tag{7}$$

and it can be found that zero reflectance can be obtained in this case with a double quarter-wave coating if the refractive indexes of the two layer satisfy the following relation:

$$\frac{n\_2}{n\_1} = \sqrt{\frac{n\_3}{n\_0}}\tag{8}$$

In Fig. **4** is shown an experimental reflectance profile of a double ARC made by MgF2 and ZnS over a flat silicon surface, having effective reflectance equal to 8.1%, good if compared to the one of a single SiN ARC, which is around 13%. However, this treatment is usually too expensive for most commercial solar cells, but it finds application in very high efficiency solar cells designed to operate in concentration. The ARC concept illustrated till now for silicon solar cells is actually

totally general and can find application on thin film cells and modules. For instance it is possible to use Silicon Nitride as ARC for thin film cells CdS/CuInSe2 [24].

**Figure 3:** Schematic of Anti Reflection Coating principle of operation with two equal thickness coating.

**Figure 4:** Experimental reflectance profile of double ARC (blue) compared with flat silicon (red).

#### **2.1. Dielectric Layers**

Dielectric layers have wide applications in photovoltaics due to their different optical and electrical properties useful for passivation, antireflection and even masking purposes.

The former two are keys to increase solar cell performances since the passivation level strongly influences the Voc, the Jsc, and in minor part the FF, while according to paragraph above the antireflection coatings are very important to drastically reduce the optical losses, and so influence the Jsc. However, not all dielectrics employed in PV are appropriate to both purpose, or suitable on all kind of processes.

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **401**

Before the wide diffusion of Silicon Nitride (SiNx) deposition *via* remote plasma [25], leading the overall silicon solar cell processing throughput at the present level, the main way to passivate the front of silicon solar cells was Silicon Oxide (SiO2) grown by thermal wet or dry oxidation in which ultra clean water vapour or oxygen respectively are entered in a quartz tube at high temperature (in the range 800-1200 °C) to oxidise the Silicon wafers. Such high temperatures are one of the drawbacks of using SiO2 in solar cells processing. Indeed metallization of cells when using silicon oxide would involve low throughput steps of lithography, since metal contacts cannot be deposited prior to high temperatures step because of the consequent lifetime degradation in the base wafer and metal contacts extra-penetration through the formed emitter [26]. Moreover since the hardness of SiO2 films, common firing through processes are not effective in this case, so an opening procedure (*via* laser or chemical processing) is mandatory. Of course silicon oxide is not suitable for thin film applications.

Sputtered or evaporated dielectric films, like Titanium Oxide (TiO2) [27] or Magnesium Fluoride (MgF2), or Zinc Sulphide (ZnS) are suitable for antireflection purpose and have been used previously in photovoltaic industry but with poor passivation quality. Plasma deposited films, like non-stoichiometric Silicon Oxide (SiOx) [28], Silicon Nitride (SiNx) also coupled with amorphous silicon (a-Si) [29] have shown both good passivation quality and optical characteristic suitable to be employed as ARC. Also they are such to be easily penetrated by common Silver screen printable paste to perform good firing through during metallization processes.

Finally, due to their low process temperatures, they are suitable also in low temperature application, like thin film [24] and heterojunction [30].

In the following paragraph we will focus on silicon nitride, which is at the present the most widely diffused dielectric employed in solar cell processing.

#### **2.2. Silicon Nitride**

The application of Silicon Nitride, deposited *via* Plasma Enhanced Chemical Vapor Deposition (PECVD), to photovoltaic represents an important step toward the enhancement of conversion efficiency, especially for c-Si based cells. Indeed its electric properties make it an excellent passivating dielectric, due to the relevant surface recombination velocity reduction, while it is a natural good candidate for Anti Reflection purpose because it is possible to continuously move the refractive index from 1.7 to 2.3 (evaluated at 633 nm), depending on deposition condition [31]. These are the main reasons to adopt this material in solar cells processing, giving a double contribute in a single step deposition. For historical reasons we report that Silicon Nitride can be deposited using high temperature process approach, but the one utilized for photovoltaic applications is deposited by the low temperature techniques (< 400 °C), like PECVD, suitable to be industrially adopted, because of low cost of ownership, high throughput and high reliability. One of the mostly interesting characteristics is the possibility of good interfacing in terms of electrical and optical properties of Silicon Nitride with different materials, like amorphous silicon [29], or Silicon Oxides [32]. Depending on deposition conditions it is possible to vary the stoichiometry composition and the Hydrogen content, that affect the fixed and mobile charge content, which are responsible of the film passivation properties. Indeed there are two main mechanisms to passivate silicon surface: the dangling bonds saturation by means of Hydrogen

atoms and a field effect due to the charge present in the film [33], better explained in the following.

Silicon Nitride can be grown in different ways that we can group in three classes:

In the two former cases the obtained film are perfectly stoichiometric (Si3N4) but, despite a long deposition process and at high temperature (circa 900 °C), very thin (< 10 nm). This low deposition rate is the main reason for abandoning this growth method toward other techniques using reaction in gas phase [34].

CVD deposition happens in a pressure and temperature controlled chamber, called reactor, because of the dissociation of gas species molecular bonding after an excitation. Depending on gas mixture, excitation source, pressure and process temperature, it can be distinguished between different CVD types [35].

Generally all these techniques use reactant gases containing Hydrogen, easily incorporated in the film. Moreover the particular deposition condition of a single solid phase from a gas one, in condition far from the thermodynamic equilibrium lead to the formation of an amorphous and nonstoichiometric material. This is the reason why, referring to silicon nitride deposited in CVD systems, they are defined as amorphous and hydrogenated silicon nitride, a-SiNx:H, (or simply SiNx) where *x* is the atomic ratio N/Si [36], equal to 1.33 for perfectly stoichiometric nitride. Both APCVD or LPCVD are based on the reaction of silane (SiH4) and ammonia (NH3) happening at the substrate surface due to the high temperature (700-1000 °C) or the low pressure in the reactor (0.01-1 Torr), with a deposition rate of around 2.5-5 Å/min [26]. Both methods are quite obsolete, due to the low growth rate and process conditions, not suitable for high throughput production in industrial environment, especially when compared to plasma assisted CVD. Similar drawbacks are typical of HWCVD, in which the gas mixture is excited by a wire kept at high temperature by current flow, which does not ensure a homogenous deposition on large area wafer surface, even with low deposition rate [37]. The main advantage of plasma assisted CVD (Plasma Enhanced CVD-PECVD) compared to the APCVD or LPCVD is the low process temperature, since the

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **403**

reaction is started by means of an external source, like an electric field in radiofrequency. In this case the electric field gives energy to electrons in the external shells of reactant gases atoms in the reactor, ionizing them, while electrons are free to travel into the gas mixture. In this way the collisions they experience with other molecules lead to new ions and radicals formation, which are then free to combine with other reactant species. During these collisions, electrons loose kinetic energy, with light emission. All this complex of ions, radicals, light and free electrons in gas phase is the so-called plasma [26]. The second strength of PECVD is the high deposition rate, which allows to grow hundreds nanometres thick film in few minutes. In summary PECVD is a fast and low temperature process, allowing the deposition of different kind of materials controling both optical and electrical properties by means of deposition parameters. PECVD systems differ by working frequencies (Low Frequency, High frequency, Very High Frequency [38], and microwave [39]), and applied plasma (direct or remote [25]). The system choice is mainly related to the application and required final properties.

As reported before, silicon Nitride properties depend strongly on the deposition method.

In APCVD and LPCVD cases the film is purely stoichiometric, very strong and poor of Hydrogen content. The material structure is made by Silicon atoms in tetrahedral coordination with Nitrogen atoms, in turn placed on a plane.

In PECVD case, instead, the ratio N/Si and the Hydrogen content are strongly dependent on the process parameters, like the precursor gases flux ratio, the RF power, the substrate temperature. Despite the non-stoichiometric and amorphous nature of the so obtained film, experiments show the tetrahedral coordinated structure for silicon and trivalent for Nitrogen is still valid [26]. The Hydrogen content increases with deposition temperature decreasing, and this is the reason why Silicon Nitride produced *via* PECVD has higher Hydrogen content with respect the ones deposited *via* APCVD or LPCVD. This phenomenon is correlated with the cross linking between Si and N atoms, because of the correspondence between the Hydrogen content and film density, in turn related to the Si-N bonding. Such a kind of behaviour can be confirmed by means of IR absorbance spectra, where depending of peaks position and intensity it can be estimated both Hydrogen content and structural properties of film, such as refractive index, related to Si-N bonds. Moreover depending on the stoichiometry of deposited material, SiNx can be classified in Silicon rich (N/Si<1.2) or Nitrogen rich (N/Si>1.4). In the former, Si-Si and Si-N bonds are definitely prevalent with respect to N-H bonds,; while in the latter, or even near the stoichiometry, Silicon can form several bonding and Nitrogen is able to establish bonds with both H and Si. N-N bonds are very difficult to be formed, even for Nitrogen rich film, due to the bonding energy (38.4 kcal/mol), lower with respect to both Si-N (50 kcal/mol) and N-H (93.4 kcal/mol). This suggests, for Nitrogen rich films, more N-H bonding than Si-H, since bonding between N atoms are less favourite, obtaining higher Hydrogen contents [40], which increases with N/Si ratio. The stoichiometric silicon nitride (Si3N4) electronic structure was calculated by Robertson [41] and shows the non-saturated bonds of trivalent Silicon (≡Si•) introduce a paramagnetic state in the middle of its bandgap (estimated to be 5.3 eV), while the defects introduced by nitrogen dangling bonds (=N•) are located just above the valence band edge. This is still valid for silicon rich SiNx (x<1.1), where trivalent silicon bonding can be saturated with different configurations, like Si3≡Si•, Si2N≡Si•, SiN2≡Si•, N3≡Si•. The bonding density is controlled by the equilibrium with weak Si-Si bonding, forming states near the band edges, while the level caused by Nitrogen

dangling bonds are not present. This happens because the high silicon content in SiNx moves the conduction and valence band edges toward midgap, reducing its value till 1.7 eV. Silicon induced defects still remain around midgap, while those due to Nitrogen disappear into Valence band. Actually the more the Silicon content, the more SiNx behaves like a-Si:H, where Nitrogen atoms act as donors, when they are less than 2%, and SiNx is, in practise, n-type a-Si [42]. From FTIR spectra it is possible to obtain an absorption spectrum as function of IR radiation wave number (expressed in cm-1) where peaks or distribution are related to the energy corresponding a particular vibration or stretching mode of a molecule, whose values are reported in literature [22, 43] and the most important indicated in Fig. **5**.

**Figure 5:** Example of FTIR spectrum for Silicon Nitride, with most important explicit peaks.

Analysing such spectra is possible to understand the material composition. In particular it can be found a relation between the presence of certain bonds with the optical and electrical material, like the refractive index and the charge inside the film [40]. The eventual peaks related to CO2 presence, around 653 and 2340 cm-1 have been reduced by software interpolation. If in silicon rich Nitride the only states at midgap are due to Silicon, different is for Nitrogen rich Silicon Nitride. In this case, also the states due to Nitrogen dangling bonds are inside the gap, and the prevalent effect is due to N3≡Si• groups, the so-called K centres of recombination.

They are diamagnetic and neutral, but introduce a large amount of instable charge in the film, since they are able to capture both electron and holes, becoming positively or negatively charged (K+ , K- ). This charge state formation can happen also after the film exposure to high energy radiation, like UV, which breaks some Si-H and SiN-H bonds around the K centre, forming new radicals and mobile charges inside the Nitride film [44]. Of course the silicon or nitrogen rich behaviour is determined by the deposition parameter, like the reactant gas flux ratio R=ΦNH3/ΦSiH4. The correspondence between deposition parameter, film properties, and atoms bonding can be appreciated in Fig. **6(a)**. We report the

**Figure 6:** a) Comparison of refractive index and N-Si3 peak intensity in FTIR spectra of Silicon Nitride deposited with different Ammonia/Silane gas flux ratio. b) comparison of refractive index and lifetime of Silicon Nitride deposited with different Ammonia/Silane gas flux ratio.

variation of refractive index (evaluated at 633 nm) as a function of R. When the film is deposited with less Silane the refractive index start to reduce, because the equilibrium between Silicon and Nitrogen atom content move their availability to form bonds with Hydrogen. Then N-Si bonding is less favourite and the film become less dense, with lower n. This behaviour is confirmed measuring IR spectra. In particular the intensity of peak located around 850 cm-1 shows the same trend of refractive index. At the same time the opposite trend in N-H and Si-H bonds is reflected in lifetime trend, as reported in Fig. **6(b)**. Indeed the higher charge content due to K centres and hydrogen is responsible of Field effect passivation on crystalline silicon: the charge inside the film can accumulate electrons near

the silicon wafer surface avoiding recombination or also inverting the silicon locally, obtaining an effective passivation. For instance in case of SiNx with high positive charge in the film, when deposited on p-type silicon, there are electrons recalled from the substrate close to the surface. This causes a band bending which moves the Fermi Level away from midgap, where the most important defects (through which electrons and holes can recombine, lowering the overall lifetime) are located [32]. This effect together with the saturation of dangling bonds at the surface defect state density can led to very high effective lifetime, in the order of milliseconds [26] both for mono and multicrystalline silicon.

**Figure 7:** Field effect: the positive charge in the SiNx recalls electron from the substrate close to the surface, avoiding recombination *via* midgap defects.

### **3. TCO AND SURFACE SCATTERING**

In thin film silicon solar cells and modules incorporating amorphous (a-Si:H) or microcrystalline (c-Si:H) silicon as absorber materials, light trapping, *i.e.* increasing the path length of incoming light, plays a decisive role for device performance. All these cells commonly have a Transparent Conductive Oxide (TCO) as top electrode deposited on low cost glass substrate. In Fig. **8** schematic sketch of the cross section of thin film solar cell is reported. The light trapping is illustrated by the arrows representing incoming/reflected and scattered sun light at different interfaces: glass/TCO, TCO/absorber layers, absorber/TCO and TCO/back reflector.

Not all device configurations include the TCO before the metal electrode but often this layer is introduced to enhance the chemical compatibility between thin films and metal electrode. In a particular kind of silicon wafer based solar cell, TCO is also applied on both sides of the cell, such as the a-Si:H/c-Si/a-Si:H heterostructure cell [45].

When TCO is applied at the cell front side, it needs high transparency in the spectral region where the solar cell is operating. Therefore refractive indexes are relevant and have to be well designed for coupling of light into the absorber materials. TCO also should have high carrier mobility and high electrical conductivity properties to reduce electrical losses. This last property is mandatory due to the poor lateral conductivity of doped emitter thin films composing the front junction. Of course the minimum requirements for the optical and electrical properties of the TCO depend on the structure and the absorber material adopted to fabricate the solar cell. Typically values of 10 /sq as a sheet resistance and a transmittance around 93-94% in the range 350-1100 nm are valuable for solar cell applications. A high transparency for visible light (wavelength from 350 nm up to 750 nm) is sufficient for amorphous silicon absorber layers [46], while for solar cells integrating c-Si:H or a-Si:H/c-Si the TCO has to be highly transparent up to the near infrared (NIR) region (1100 nm). This reflects in certain restrictions on the carrier density n of the TCO

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **407**

material. Indeed higher the free carrier absorption, lower is the IR transmittance [47, 48]. However the success of the TCO films is mainly determined by the physical-chemical properties for the growth of the thin film material on top of it. For example it has to be inert to hydrogen-rich plasmas or act as a good nucleation layer for growth of nano-or microcrystalline material. A brief summary of the relevant properties of the most common TCO for photovoltaic applications, such as ZnO, SnO2 and ITO is listed in the following section.

**Figure 8:** Schematic cross section (not in scale) of thin film cell with rough interfaces. The light trapping is remarked by the arrows representing impinging and scattered sun light at the rough interfaces.

## **3.1. ZnO**

ZnO films is n-type semiconductor due to oxygen deficiencies with an energy gap of 3.3eV, commonly doped by adding aluminium, gallium, boron or indium [49].

In Fig. **9** resistivity as a function of impurities is reported on the base of [50]. In the same work it is shown that a film with low resistivity (5.10-4 cm) and high optical transmission (greater than 85% between 400 and 800 nm) can be produced on low-temperature substrates with a relatively high deposition rate in particular RF magnetron sputtering configuration. The sheet resistance and Hall mobility of the film are 10 /square and 120 cm2 /Vs respectively [50].

**Figure 9:** Dependence of ZnO resistivity () on impurity/Zn ratio in the target. Film with B(▲), Ga (∆), Al(●), In(□) have been experimented [50].

Theoretical work [51] has identified also hydrogen as a possible n-type dopant although an enhancement of conductivity by hydrogen treatment has been known for long time [52].

Various methods are commonly used for TCO preparation. The methods involve spray pyrolysis, vapour and liquid phase or sol–gel deposition, pulsed laser deposition and sputtering. Zinc oxide is often prepared by CVD processes (using, *e.g.*, Zn(C2H5)2 as a precursor gas) or sputter deposition [53]. Substrate temperatures differ quite significantly in these two deposition techniques. Indeed, for ZnO prepared by CVD, temperatures near 150 °C are applied [53, 54], while for ZnO prepared by sputtering, typical process temperatures are near 300 °C.

Recently doped ZnO has been established as a successful alternative to the ZnO also on large area module as demonstrated by Oerlikon [55]. Examples are boron doped ZnO prepared by lowpressure chemical vapour deposition (LPCVD) [56, 57] or aluminium-doped films deposited by expanding thermal plasma CVD [58]. A very helpful approach is the use of magnetron sputtering together with a post-deposition chemical etching step. The sputtering process leads to highly conductive and transparent but smooth ZnO:Al films. Typical sheet resistance values of 5–8 /sq for film of 650 nm thicknesses and absorption of only 5–6% in the wavelength range from 350 nm to 1100 nm proof the high quality of these films. A surface texturing can be performed by a simple chemical etching step in diluted hydrochloric acid (typically 0.5% HCl) to improve the light scattering as well as the light trapping of the ZnO film. This etch yields a textured surface which can be adjusted to give optimal light scattering over a wide wavelength range. Fig. **10** shows scanning electron microscope (SEM) images of an (a) as deposited and (b) etched ZnO:Al film. Sputtered ZnO:Al is initially smooth. With increasing etch time, the roughness of the surface increases from 7 nm to 162 nm [49].

ZnO electrical properties are sensitive to the working pressure and the deposition temperature as demonstrated in Fig. **11** [48].

**Figure 10:** SEM images of RF sputtered ZnO:Al surface before (a) and after wet chemical texturing (b) [49].

It should be noted that film properties depend also strongly on film thickness. In Fig. **12** the Hall mobility, carrier densities and conductivities of Al doped ZnO films are reported according [48] as a function of deposited thicknesses and for different doping density related to the concentration of Al2O3 in the target (referred as TAC).

**Figure 11:** (a) resistivity (), (b) carrier concentration (N) and mobility (), of Al:ZnO layers as a function of working pressure during film deposition [48].

**Figure 12:** (a) Resistivity (), (b) mobility (), (c) carrier concentration (N) of Al:ZnO layers as a function of film thickness [48].

Also the optical properties can be affected by the deposition temperature as evident from the transmittance spectra of ZnO films deposited at different temperature and different doping densities reported in Fig. **13**.

Electronic transport parameters such as sheet resistance *Rsh*, resistivity , carrier concentration *N*, mobility , and the optical gap for ZnO:Al films deposited with different target doping concentration TAC and deposition temperature are listed in Table **1**.

**Table 1:** Thickness; Sheet resistance, Rsh; Resistivity, ; Carrier concentration, N; Mobility, and Optical gap, E Gap for ZnO:Al films deposited with different target doping concentration ( TAC) and deposition temperature, T.

When viewing the electrical parameters it can be seen that the lowest resistivities and the highest mobilities have not changed in the last 15 years, despite a lot of technological effort in this period [53]. The lowest resistivities reported are in the range of 1.4 to 2×10−4cm, independently of the deposition method. Carrier concentration of 5.1020 cm−3 and mobility of about 40 cm2 V−1 s −1 , can be considered as an upper limit. It seems reasonable to suppose, that resistivities less than 1.10−4 cm for doped polycrystalline zinc oxide films are hard to achieve, due to a general limitation by ionized impurity scattering. This assumption is further supported by the fact that heavily n-doped single crystalline silicon also exhibits mobilities in the same range as polycrystalline ZnO films [59].

**Figure 13:** Transmittance and reflectance of different doping films (Ts =100°C and film thickness d=600-725 nm) deposited from different percentage Al content in the target [48].

Conducting p-type ZnO has recently been realized [60]. Even if the conductivity still remains on open request ( around 300 cm), a promising way to enhance the light trapping in thin film cell using ZnO film, concerns the formation of surface nanostructures. They exhibit a high crystalline quality. Fig. **14(a)** shows a tilted SEM image of the ZnO nanorods synthesized on the fluorine thin oxide substrate by chemical solution deposition (CSD). The SEM image in the inset of Fig. **14(a)** shows a cross-sectional view of the ZnO nanorods. The length and average diameter of the ZnO nanorods were about 1.5 μm and 80–90 nm, respectively. Each ZnO nanorod has a hexagonal shape. The SEM images of Al n-doped p-type ZnO thin film coated ZnO nanorods are depicted in Fig. **14(b)**. The thickness of the ZnO thin film coating the nanorod is about 200 nm [61].

Very helpful ZnO nanorods for antireflection coating in solar cell application is also reported in ref [62] where the authors achieved a weighted global reflectance of 6.6% on silicon surface.

## **3.2. SnO2**

As a front transparent conducting oxide (TCO) layer for thin-film solar cells, ZnO:Al has a lot of advantages over the conventional SnO2:F, *e.g.*, high stability under the hydrogen plasma environment and high electro-optical properties [63]. In turn, due to the higher contact potential of ZnO:Al/p a-SiC:H compared with SnO2:F/p a-SiC:H, the contact properties of n-type ZnO:Al with adjacent p-type a-SiC:H can be limiting factors for a-Si:H solar cell performances, since the a-Si:H solar cells prepared on ZnO:Al suffer from reduced fill factor and Voc [64].

**Figure 14:** (a) SEM image of ZnO nanorod synthesized on F:SnO2 substrate. Inset shows the cross sention view. (b) Tilted SEM image of the p-type ZnO thin film coated ZnO nanorod [61].

Fluorine-doped tin oxide films (SnO2:F), which fulfil these requirements up to a large extent, have been developed by Asahi Glass (Asahi Type U, [65]) by chemical vapour deposition CVD process using SnCl4 precursor gas [56, 57]. It is also the most commonly used TCO material for thin film Si solar large area module manufacturing [65].

Accurate optical constants (n and k) of textured SnO2 films have been reported in the range of 0.85–4.6 eV using spectroscopic ellipsometry on Ashai U type SnO2 on glass substrate [66]. In Fig. **15** is reported the transmission spectrum of the SnO2 film deposited on glass substrate.

In Fig. **16** are reported the refractive index n as well as the extinction coefficient k of the textured SnO2 on glass measured at different angles. For comparison the n and k of SnO2 flat surface are also reported.

SnO2:F can also be fabricated by spray pyrolysis technique as referred in [67]. In that experience SnO2:F films were deposited from a solution comprising 23 ml of SnCl4, 7 ml of H2O, and 920 ml of methanol solution 6. Fluorine was added in the form of HF, instead of NH4F, to give the solution a doping level of F/Sn in the range of 0–180 atomic per cent.

**Figure 15:** Measurements and fit of the transmission measured on SnO2 deposited on glass substrate [66].

Deposition was performed on borosilicate glass substrates at a temperature of 490 °C-525 °C and from a distance of approximately 18 cm using a spray bottle with a nozzle diameter of 0.5 mm. The carrier gas was nitrogen and the spray rate was 12 ml/min. , N as a function of doping level expressed as F/Sn atomic percentage are reported in Fig. **17**.

**Figure 16:** Refractive index n and extinction coefficient k of the textured SnO2 on glass measured at different angles. For comparison dashed line refers to n and k of SnO2 flat surface [66].

The transmission spectra of doped (n = 7.2.1020cm-3, = 4.7 10-4cm thickness = 0.25 m) and undoped (n = 3.2.1020cm-3, = 63.4 10-4cm thickness = 0.28 m) SnO2 films are reported in Fig. **18**.

## **3.3. ITO**

Basically the ITO is a degenerate n-type doped semiconductor that can be fabricate by thermal evaporation [68] or by sputtering system [69, 70]. Here we briefly report on the ITO films as fabricated by the authors for amorphous crystalline heterojunction solar cell solar cell applications.

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **413**

Concerning the first technique, the ITO was deposited using a Balzers BA510 electron gun system starting from a self-made target of Indium and Tin alloy. Films of 150 nm thickness have been deposited on Corning 7059 in oxygen atmosphere setting a chamber pressure of 2·10-4 mbar. Temperature was measured directly on the substrate using a K thermocouple. Sn/(Sn+In) (wt%) target ratio and substrate temperature were varied to obtain minimum resistivity and best transmittance in the visible range.

**Figure 17:** , N as a function of doping level expressed as F/Sn atomic percentage [67].

**Figure 18:** The transmission spectra of doped and undoped SnO2 films [67].

The variation of resistivity and averaged transmittance from 400 to 800 nm as a function of Sn/(Sn+In) ratio for a fixed substrate temperature are reported in Fig. **19**. The variation of resistivity and averaged transmittance from 400 to 800 nm as a function of substrate temperature for a fixed Sn/(Sn+In) ratio are reported in Fig. **20**.

The ITO films useful for a-Si:H/c-Si heterostructure solar cells [71, 72] have been deposited using a low temperature RF 13.56 MHz magnetron sputtering. In Fig. **21** are reported, as a function of deposition temperature, the values of resistivity and effective transmittance from 350 nm to 1200 nm defined as:

$$T\_{\eta^{\text{c}}} = \underbrace{\int\_{10^{\text{c}}}^{10^{\text{c}}} T(\lambda) \cdot AM \, 1.5G(\lambda)d\lambda}\_{\text{(c)}^{\text{(L)}}} \underbrace{\int\_{10^{\text{c}}}^{10^{\text{c}}} T(\lambda) \cdot AM \, 1.5G(\lambda)d\lambda}\_{\text{(L)}^{\text{(L)}}} \tag{9}$$

**Figure 19:** Resistivity and optical averaged transmittance (400-800 nm) of ITO *vs.* percent of Tin in the Tin+Indium solid source for films deposited at 260°C of temperature substrate [68].

Sn/(Sn+In) (wt %)

**Figure 20:** Resistivity and optical averaged transmittance (400-800 nm) of ITO *vs.* substrate temperature of deposition for fixed target composition (60% Sn content) [68].

where T() is the transmittance and AM1.5G() is the number of photons per unit wavelength ().

**Figure 21:** Sputtered ITO resistivity and transmittance as a function of deposition temperature.

**Figure 22:** Sputtered ITO resistivity and transmittance as a function of oxygen flux at a deposition temperature of 80°C.

The deposition process parameters have been optimized in order to obtain a good compromise between the film conductivity and transmittance. In particular, the sputtering conditions have been fixed as follows: the RF power to 200 W, the process pressure to 2.25 mTorr, the Ar content to 25 sccm. It is worth to note that no oxygen has been used in order to reduce the film resistivity. For comparison the effective transmittance and resistivity as a function of oxygen flux during the sputtering process at a fixed temperature of 80°C are reported in Fig. **22**.

## **3.4. Texturing**

Basically light trapping is an interesting technique to increase absorption of the incident light in the active layers of solar cells. This is particularly true in all thin film cells where the absorber layer thickness is chosen on the base of the orthogonal light absorption. To enhance the optical path of the refracted light impinging onto the cell, a texturing technique is commonly adopted producing roughness on the window layer of the device. Moreover to further enhance the internal absorption a rear side textured reflector is used, thus producing several light round trip within the absorber layer. These concepts can be successfully extended also to solar cell involving thicker absorber layer, such as silicon wafer based cell, even more when wafer thinning is mandatory to reduce the PV cost. In this paragraph we focus on texturing techniques used in thin film and in c-Si based solar cell also going toward thinner silicon wafer.

Thin film cells are commonly based on different kind of architecture depending on the choice of mechanical support as substrate. If a hard substrate such a glass is used, then a top down configuration is adopted in the cell manufacturing, thus the glass is the cell window layer [73]. On the other hand if a flexible substrate is preferred such as stainless steel, then the cell fabrication process has a bottom up development during the growth [74]. These are the most common configurations, even if mixed approach are possible such as flexible transparent plastic foil [75] or hard metallic substrates. Except for the CdTe and CIS/CIGS cells, in which the TCO is used just for the front side contact, the other device commonly use TCO for both front and rear side contacts. In Fig. **23** a schematic overview of the different configuration is reported. To enhance the cell absorption by the texturing technique, two strategies can be adopted and eventually merged: TCO texturing and substrate texturing. Both introduce light scattering at the rough interfaces in the solar cell. In this way the average length of the optical light path in a layer is increased and the light absorption enhanced. Scattering at rough interfaces not only leads to enhancement of the absorption in the active layer of the cell but also influences the absorption in all other layers.

**Figure 23:** Different thin film cell configurations.

Now we will consider the scattering promoted by a rough surface of TCO deposited on flat glass surface as depicted in Fig. **24**.

**Figure 24:** Light scattering at a thin film with one roughly textured surface deposited on glass surface.

For an optimal optical design of thin film solar cells it is important to know the light absorption profile in the whole solar cell. Since it is not easy to determine this profile experimentally, computer modeling has often reported in literature to study the absorption in thin film solar cells [76-78]. The model is based on the first order Born approximation [79] and on Fraunhofer scattering. A rough interface can be described taking into account the root mean square (rms) roughness,rms and the scattering data including the haze parameter and the angular distribution of diffuse light. The rms is defined as follows:

$$
\sigma\_{r\text{nu}} = \sqrt{\frac{1}{N} \sum\_{i=1}^{N} \left(\eta\_i - \bar{\eta}\right)^2} \tag{10}
$$

Where is the height of the i-th scanned point, the average height of the scanned point and N is the total number of point collected during an atomic force microscopy (AFM) scan.

The haze for transmitted light normally incident to the glass side and crossing a rough surface as depicted in Fig. **24** is defined:

$$\mathbf{H}\_{\rm T} = \mathbf{T}\_{\rm diff} / \mathbf{T}\_{\rm tot} \tag{11}$$

where Ttot is the total and Tdif is the diffused optical transmittance measured by a photospectrometer equipped with an integrating sphere. The total and diffused transmittance of a fluorine doped tin oxide (FTO) SnO2 glass coated, deposited using atmospheric pressure chemical vapour deposition APCVD, commonly available on market as Ashai U-type [65], having around 40 nm, is reported in Fig. **25**.

To simulate the haze of TCO it is possible to use the following equation:

$$H\_T(\mathcal{Z}) = 1 - \exp\left[ - \left( 4\pi\sigma\_{mu}C\_T \left| n\_1(\mathcal{Z})\cos\theta\_{m} - n\_2(\mathcal{Z})\cos\theta\_{out} \right| \mathcal{\lambda} \right)^{\rho} \right] \tag{12}$$

are fitting parameters. While a square dependence ( = 2) is reported by the scalar scattering theory [80], other research group found better agreement for large wavelength as cubic dependence [76].

The meaning of CT is quite controversial: it is dependent from and it is strongly needed to have reasonable results. Physically might be related to the shadowing effect of scattered light [81]. Power factors larger than 3 for ZnO:Al layers textured by post-etching in HCl is also reported in [82, 83].

**Figure 25:** Total and diffused transmittance together with haze of FTO Asahi U-type.

**Figure 26:** AFM scan image over a 3x3 m2 of F:SnO2 Ashai U-type.

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **419**

Moreover is not easy to evaluate the rms because, if atomic force microscopy (AFM) is used, the rms can depend on the scanned area, becoming higher for larger scanned area. This would indicate that the roughness varies more than the range of one micron as a typical AFM scan. Therefore large area should be investigated with AFM to estimate reasonable rms values.

After deposition of absorber layer on rough TCO surface a smoothening occurs that reduces the rms. Indeed, in case of P-I-N amorphous silicon solar cell deposited onto Asahi U-type covered corning glass, while the very short p-type doped layer almost does not affect the TCO roughness, after 300 nm intrinsic layer deposition the rms reduces to 33 nm [77]. An AFM scan image over a 3x3 m2 of F:SnO2, Ashai U-type is reported in Fig. **26** [83]. A SEM picture remarking the pyramidal surface of this textured TCO is reported in Fig. **27** [76].

**Figure 27:** The SEM picture of the rough surface of the glass/SnO2 Asahi U-type substrate [76].

A haze comparison of three different ZnO textured surfaces in case of ZnO/air and ZnO/c-Si interfaces are respectively reported in Fig. **28(d), (e)**, together with the simulations obtained using equation (12). The thicknesses, the rms and the shape of these Al doped TCOs are listed in Table **2**. Moreover the shape of the three different Al ZnO textured surfaces are depicted in SEM images reported in Fig. **28(a), (b), (c)** [84].

**Table 2:** Parameters related to the Al doped textured ZnO.

**Figure 28:** SEM pictures (a), (b, (c) of typical type-A,-B and-C ZnO layers respectively. Experimental (symbols) and calculated (lines) haze in transmission in air (d) and at the ZnO/Si interface for the type-A,-B and-C front ZnO layers (e) [84].

Basing on the experience reported in [77], now we report about the use of model related to the equation (12) to evaluate the reflectance and transmittance of a Ashai U-type F:SnO2 used for p-in a-Si:H solar cell.

In case of rough interface, having rms smaller than /n of the medium incidence, the scalar scattering theory is applicable to determine the specular reflectance of an interface. The total reflectance Rs can be modelled using the following equation reported in [77]:

$$R\_S(\mathcal{A}) = R\_0(\mathcal{A}) \exp\left[ - \left( 4\pi \sigma\_{\rm rms} n\_0 / \lambda \right)^2 \right] \tag{13}$$

where R0 is the reflectance of a flat surface. Assuming that:

$$R\_0(\mathcal{Z}) = R\_i(\mathcal{Z}) + R\_d(\mathcal{Z}) \tag{14}$$

where *Rd()* is the diffused reflectance, we obtain

$$R\_d(\mathcal{A}) = R\_0(\mathcal{A}) \left( 1 - \exp\left[ -\left( 4\pi \sigma\_{mu} n\_0 / \lambda \right)^2 \right] \right) \tag{15}$$

Fig. **29(a** and **b)** show the ratio *Rd/R0* as calculated from equation (15) for light incident at a rough interface through Asahi U-type TCO and n-type a-Si:H, respectively, for several values of rms. It

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **421**

is important to note in Fig. **29** that the back contact interface acts as a nearly perfect diffuser for the reflected light in the wavelength range of interest when rmsis larger than 30 nm [77].

**Figure 29:** The *Rd*/*R*0 ratio as a function of wavelength for several values of the rms roughness of the rough interface: a) Asahi U-type TCO material and b) *n*-type *a*-Si:H [77].

To evaluate the diffused transmittance Td() of the rough Ashai U-type TCO it is possible to use a relation similar to equation (15):

$$T\_d(\mathcal{A}) = T\_0(\mathcal{A}) \left( 1 - \exp \left[ -\left( 4 \pi \sigma\_{ms} C \left| n\_0(\mathcal{A}) - n\_l(\mathcal{A}) \right| / \, \lambda \right)^3 \right] \right) \tag{16}$$

where C depends on the two media and in particular when |n0-n1| increases C tends to 1. Fig. **30** shows the ratio of Td to T0 for TCO/p and p/i interfaces for different values of rms assuming C = 1.

By using the modelled transmittance and reflectance and the values of rms at each interfaces it is possible to evaluate how the interfaces within the TCO/p/i/n/Ag structure affect the Quantum Efficiency of the solar cell. Indeed a Quantum Efficiency can be seen as the sum of optical and electrical contributes arising from each scattering interface and each layer respectively.

$$\underline{QE}(\lambda) = \sum\_{\text{hyp}} \underline{QE}\_{op}(\lambda) \underline{\eta}(\lambda) \underline{QE}\_{el}(\lambda) \tag{17}$$

QEop is the optical quantum efficiency measuring the probability of a photon to be absorbed; () is the carriers generation per photons; QEel is the collection probability.

**Figure 30:** The *Td/T*0 ratio as function of wavelength for several values of rms for: a) the TCO/*p* interface and b) the *p*/*i* interface [77].

The texturing effects on front and rear contact on experimental QE of four different TCO/p/i/n/Ag cells are reported in Fig. **31** as points. The cells have different intrinsic layer thicknesses. For each of them four different simulations are reported in the same figure as lines. Each of them refers to: (1) all interfaces are flat; (2) only TCO/p is a rough interface; (3) only n/Ag is a rough interface; (4) all interfaces are rough. For each case it is possible to evaluate the short circuit current (Jsc) integrating the QE values over the wavelength range from 350 nm to 900 nm and taking into account the photon number per wavelength of sunlight AM1.5G spectrum. This value is also reported in Fig. **31** for each simulation as Jph. The simulation clarifies that the scattering enhances the absorption in the intrinsic layer leading to a gain up to 2.3 mA/cm2 . This enhancement is slightly dependent on the intrinsic layer thicknesses. The simulation also states that the roughness close to the rear contact has a strong effect in the spectral region of wavelength higher than 650 nm and has no effect on front side of the cell then does not affect the QE in the spectral range from 350 nm to 550 nm. Moreover the front rough interface strongly affects the QE in the spectral range between 450 nm and 650 nm slight, while introduces a slight decrease of the QE below 450 nm.

Introducing a TCO layer with rough surface in between n/Ag interface it is possible to further enhance the QE in the spectral range above 650 nm, leading to a Jsc improvement up to 2.9 mA/cm2 .

As expressed before, the angular distribution of diffused light is a relevant parameter to evaluate the scattered light in transmission IT and then to extract information on the surface roughness of the investigated medium. A very helpful tool to determine IT is the Angular reflectance transmittance Analyzer (ARTA) [86] depicted in Fig. **31**. This instrument uses monochromatic

light coming from a spectrophotometer. The detector is located in a little integrating sphere and can be moved along the sphere for an angle between-165° and 165°. The sample under test is placed in the centre of the sphere and can be rotated with respect to the incident light from 0° (normal incidence) to 90° (parallel incidence). At normal incidence the reflected light cannot be collected due to restriction of the detector

**Figure 31:** Experimental (■ points) and simulated (lines) QE of the four single junction *a*-Si:H solar cells each with a different thickness of the intrinsic layer: (a) 150 nm, (b) 300 nm, (c) 450 nm, and (d) 600 nm. The simulated QE curves refer to: (□) all flat interfaces; (∆) only the front interfaces TCO/*p* and *p*/*i* are rough; (o) only the back interfaces (*i*/*n* and *n*/metal) are rough; (full line) all interfaces are rough. The rms roughness of the front and rear interface of the cell and the Jph values are reported for each cell [77].

The absorbance A of the sample is related to the incoming light intensity I0 as follows:

010 *A I I* **(18)**

In this way the IT can be measured over a wide range of wavelengths.

Several textured substrates have been evaluated as reported in Table **3** where the surface structure and the rms values are listed. The FTO has been already described.

**Figure 32:** Schematic cross section of the ARTA [85].

**Table 3:** Morphology and rms of different textured TCOs.

**Figure 33:** AFM scan image of Al:ZnO etched surface.

The Al:ZnO (AZO) has been deposited in condition so that it resulted flat after the deposition and then has been etched in 0.5% HCl bath for 40 s to form a crater like surface morphology as depicted by AFM scan image reported in Fig. **33** [87].

#### *Light Trapping in Thin Silicon Solar Cells Silicon Based Thin Film Solar Cells* **425**

Boron depod ZnO (BZO) has been deposited by LP-CVD. The transmitted light angular distributions of samples listed in Table **4** are reported in Fig. **34(a),(b)** for 450 nm and 700 nm monochromatic incident light respectively. To compare these intensities each other they are normalized to the total transmittance Ttot(). From the measurements reported in Fig. **34** it is evident that at 700 nm the higher the rms, the stronger the scattering; moreover the AZO sample shows a stronger scattering at smaller angles than BZO. On the contrary, for both wavelengths, the specular part of the transmittance decreases with increasing rms.

**Figure 34:** Scattered light angular distribution at = 450 nm (a) and 700 nm (b) for the samples listed in Table **3** [85].

Now we report about an interesting study on the scattering properties of ZnO useful for thin film solar cells [87]. They have used the most common way to fabricate ZnO films by 13.56MHz rf magnetron sputtering system using a ZnO:Al2O3 target. Four different kinds of ZnO films have been prepared on 1737 Corning glass by varying the Al2O3 weight percentage (TCA) of the target as: 0.2, 0.5, 1 and 2%. 0.3 Pa deposition pressure has been used for three kinds of film, while 1 Pa has been adopted in case of 2%. Different films have been deposited at different temperature in the range between 60 and 490 °C. After deposition a wet chemical etch in 0.5% HCl bath has been performed on top of the film initially smooth to texture the surface. By varying the deposition conditions the ZnO films can assume different optoelectronic properties. Briefly we summarize the most relevant, such as carrier concentrations-mobilities as a function of deposition temperature for different TAC, reported in Fig. **35**, and transmission-absorption in the wavelength range between 300 nm and 2500 nm, reported in Fig. **36**.

**Figure 35:** Carrier concentration and mobility of Al:ZnO films as a function of deposition temperature for different TAC. Reprinted with permission from [87]. Copyright (2007), American Institute of Physics.