M OTIVATION - 非晶矽 / 微晶矽疊層薄膜太陽能電池之模擬與微晶矽結晶率之探討

CHAPTER 1 INTRODUCTION

1.4 M OTIVATION

The solar spectrum consists of energies of wide range wavelengths. For efficient use of solar energy, multi-junction solar cells, such as double-junction (tandem) or triple-junction cells are used. The tandem solar cell can be fabricated with only silicon active layers and therefore makes it a preferable solar cell techniques. A tandem solar cell consists of a top and a bottom cell. A high bandgap material (a-Si:H) is used for the top cell, less thermalization losses occur as a result of absorption of high energy photons. The longer wavelength photons, which are not absorbed in the top cell, get absorbed in the bottom cell which consists of a lower bandgap material (µc-Si:H). Based on the review from the literature, we set up optical and electrical simulations for single-junction and tandem solar cell in order to compare with the realistic device and analyze the details of physics. Improvements regarding optical and electrical properties of the tandem solar cell are investigated to promote the performance of the solar cell.

From the experimental aspects, intrinsic hydrogenated microcrystalline silicon (µc-Si:H) has been shown to be a very promising new photovoltaic material for thin-film solar cells.

However, the composition of µc-Si:H is complicated that it is hard to control the crystallinity.

In this study we also try to control the crystallinity of µc-Si:H by modulating the growth parameters of PECVD to reach the crystallinity of 50% which is believed to be the optimized value reported by literatures.

Chapter 2 Literature Review

2.1 Silicon Thin Film Solar Cell

Several thin-film technologies have utilized silicon for the active material. Sticking to silicon instead of other semiconductors represents a number of advantages. Silicon is nontoxic, which makes it easily accepted by the public. Moreover, silicon is abundantly available in the earth crust, so that its availability (at least in its raw form) will never be an issue. Finally, silicon solar cell technologies can build further upon the extensive know-how accumulated over the years in the IC industry (for crystalline silicon) and the display industry (for amorphous and microcrystalline silicon). The term “thin-film silicon” is in fact quite broad. It covers a wide range of technologies from amorphous to microcrystalline silicon with thickness ranged from 0.1 to 5 μm. This section intends to give an overview of the solar cell technology containing amorphous and microcrystalline silicon thin film.

2.1.1 Single-Junction Amorphous Silicon Solar Cells

Amorphous silicon is usually deposited using the plasma-enhanced chemical vapour deposition (PECVD) technique, and the gas silane (SiH4) is mostly used as precursor. As a result of decomposition, surface adsorption and surface reactions, a network of Silicon atoms is formed on the substrate, mostly glass or a metal foil. The deposition temperature ranges typically between 180^◦C and 280^◦C. As a result of the deposition mechanism, amorphous silicon contains a large concentration of hydrogen atoms (~10%). Hydrogen is, in fact, crucial for the material’s electronic properties, while unhydrogenated amorphous silicon is of no use for devices. Therefore, the material one usually refers to using the words “amorphous silicon”

is in fact hydrogenated amorphous silicon (a-Si:H). Amorphous silicon is a material that features short-range order but lacks long-range order. As in crystalline silicon, each silicon

atom is mostly fourfold coordinated, but the bond lengths and angles between the bonds show a wide variation. This structure has a strong impact on the electronic structure of the material.

Because the structure is no longer periodic, the strict conservation of momentum does not hold. As a result, instead of the indirect bandgap of crystalline Si, it basically has a direct bandgap. Therefore the absorption coefficient in a-Si:H (1.7 to 1.9 eV) is much higher than that in crystalline Silicon. However, a-Si:H suffer from degradation upon exposure to sunlight [5]. This phenomenon, called the Staebler-Wronski effect, causes a large increase in defect density (strong decrease in excess carrier lifetime), and is reversible upon annealing at temperatures above 150 ^oC. From values between 10¹⁵ and 10¹⁶ cm⁻³ in annealed state, defect densities increase to ~2 x 10¹⁷ cm^-3 in light-soaked state. The metastable defects are believed to be dangling bonds formed by breaking weak bonds in the random network. The defect densities mentioned are valid for intrinsic amorphous Silicon. Doped amorphous Silicon, obtained by adding diborane to the gas flow for p-type material and phosphine for n-type, contains much more defect than intrinsic a-Si:H (several orders of magnitude higher).

Therefore, only intrinsic a-Si:H can be used as an absorber material. Fig. 2-1 shows schematically the results obtained by W. E. Spear et al. [6] by plotting the values of dark conductivity and dark conductivity activation energy Eσ against gas phase doping ratio. Also plotted is the estimated position of the resulting Fermi level EF, obtained by taking the Eσ

values and correcting for the so-called statistical shift. The key parameters for amorphous Silicon are its dark conductivity, photoconductivity, and its mobility lifetime product. The conductivity of intrinsic a-Si:H in the dark is extremely low (< 10⁻¹⁰ S/cm) because of the low mobility (a-Si:H layer is around 20 cm²/Vs at best), the large bandgap, and the fact that charge carriers at low concentration are trapped at defects. Under illumination, many of the defects get filled with photogenerated carriers and are saturated. As a result, many more charge carriers are available for charge transport, and the conductivity is many orders of magnitude higher than that in the dark. The photoersponse, defined as the ratio of the illuminated

conductivity to the dark conductivity, is a good indication for the suitability of the material for devices, and should be larger than 10⁵.

Fig. 2-1 The dark conductivity σ, activation energy Eσ, and estimated position of Fermi level EF for a-Si:H, as a function of gas phase doping ratio NPH3/NSiH4 (for n-type layers) and NB2H6/NSiH4 (for p-type layers). Eσ* is the estimated ‘true’ distance between band edge (Ec, Ev) and the Fermi level EF, where the statistical shift Es has also been taken into consideration for n-type layers, assuming a constant defect density of 10¹⁶ /cm²eV.

For p-type layers, an identical correction Es has been assumed. In the graph, the equivalent bandgap of a-Si:H, or the ‘mobility gap’, istaken to be 1.7 eV.

The mobility lifetime (μτ) product is the crucial parameter for the transport properties of excess charge carriers in the layer, and in device grade, amorphous Silicon is larger than 10⁻⁷ cm²/V. As a result of the low carrier mobility and low lifetime, collection cannot take place through diffusion. A strong drift field is an absolute requirement. Overall, it is not advisable to use amorphous silicon p–n-type diodes as solar cell structure for three reasons. (1) The doping

capability of a-Si:H is rather poor, the Fermi-level can be pushed only half way towards the conduction and valence band edges, even with heavy doping—this can be seen in Fig. 2-1. (2) Doping has a detrimental effect on a-Si:H layer quality, because it leads to the creation of many additional silicon dangling bonds, which are the main recombination centers in this material. (3) In a classical p–n-type solar cell, carrier collection is obtained by minority carrier diffusion within the p- and n-layers. Diffusion lengths in crystalline silicon wafers are sufficiently high (over 200 μm), to ensure carrier collection over the whole useful range of the solar cell thickness where significant optical absorption takes place. In a-Si:H layers, minority carrier diffusion lengths are extremely small (around 0.1 μm), and it becomes impossible for collection of photogenerated carriers on diffusion. Because of these three reasons, p–i–n diodes are always used for a-Si:H solar cells.

When modeling thin-film silicon devices, it is important to take into account the electronic structures of a-Si:H and μc-Si:H. The spatial disorder in the atomic structure of a-Si:H results in a continuous density of states (DOS) in the band gap with no well-defined conduction-band (CB) and valence-band (VB) edges. When considering the transport properties of charge carriers in a-Si:H, we have to distinguish between the extended states and the localized states in the DOS distribution as sketched in Fig. 2-2. The localized states within the mobility gap strongly influence the trapping and recombination processes; therefore, the trapped charge in the localized states cannot be ignored, as is often the case in the modeling of crystalline semiconductor devices. The localized states in the mobility gap of a-Si:H are represented by the CB and VB tail states and the defect states. These states are different in nature. The tail states behave like acceptor-like states (CB-tail states) or donor-like states (VB-tail states), and their density is described by an exponential decay into the mobility gap.

The most common defect in a-Si:H is a dangling bond. A dangling bond can be in three charge states: positive (D⁺); neutral (D⁰); and negative (D⁻). An imperfection with three possible charge states acts to a good approximation like a pair of two imperfections consisting

of a donor-like state (DB^+/0) and an acceptor-like state (DB^0/−), and is therefore represented by two so-called transition-energy levels E^+/0 and E^0/− in the band gap. The continuous density of defect states is represented by two equal (Gaussian) distributions located around the middle gap. The corresponding pair of defect energy states is separated by the correlation energy. The different nature of the localized states in a-Si:H requires different approaches for the calculation of recombination-generation (R-G) statistics through these states. The models that are commonly used to describe the localized states in a-Si:H. The energy states in the bandgap act as trapping and recombination centers and therefore strongly affect many electronic properties of a-Si:H and the performance of a-Si:H devices. In contrast to crystalline semiconductors, in which the recombination process is typically dominated by a single energy level in the bandgap, in a-Si:H, contributions from all bandgap states to the recombination–generation (R–G) rate are included. In order to model the recombination process through the single level states, such as localized tail states, Shockley-Read-Hall R–G statistics is applied for the amphoteric defect states.

Fig. 2-2 Density of states of a-Si:H [4].

2.1.2 Single-Junction Microcrystalline Silicon Solar Cells

Microcrystalline is a mixed phase material, containing a crystalline silicon fraction and an amorphous silicon fraction. The crystallites are generally only a few nanometers to a few tens of nanometers in diameters, and are present in “bunches” or“conglomerates” in the layers. These conglomerates are much larger than the crystallites themselves, up to a micron or even larger. Because the crystallites are in the nanometer range, microcrystalline silicon is often referred to as “nanocrystalline silicon”. The two names are nowadays used interchangeably. Like amorphous silicon, microcrystalline silicon contains a lot of hydrogen (several percents), which is incorporated in situ during deposition and ensures passivation of most defects in the layers. The term “microcrystalline silicon” covers, in fact, a whole range of materials, ranging from amorphous silicon with a few percents of crystalline phase to a material with only a few percents of amorphous silicon. The properties of the materials at the two extremes are quite different, and one has to pay attention not to generalize properties that are, in fact, only valid for a limited range of crystallinity. In practice, the best devices are obtained with material close to the edge between microcrystalline and amorphous Silicon, so most recent papers refer to this type of material, which contains a large amorphous fraction.

Like amorphous Si, microcrystalline Silicon is typically obtained by PECVD at low temperature (between 100 and 300^◦C). Usually, a large hydrogen flow is added, which results in microcrystalline Si instead of amorphous Si (“hydrogen dilution”). The very high frequency (between 30 and 300 MHz) leads to a softer ion bombardment, which is more favorable to microcrystalline Si formation, and, at the same time, allows relatively high growth rates. It is, however, possible to obtain excellent results using the standard RF PECVD technique at 13.56 MHz, provided the right parameters in terms of pressure and gas flow are selected. The conditions used are the so-called “high pressure depletion” (HPD) conditions, where the relatively high pressure (~10 Torr) ensures that ions lose a lot of their energy before reaching the surface. It is important to ensure a high hydrogen content in the plasma in HPD

regime, which is usually obtained by a high hydrogen flow. The combination of Very High Frequency and HPD conditions has led to excellent solar cells. An important aspect for all deposition techniques is the need to control the crystallinity profile of the microcrystalline layer throughout the active layer. As the material is formed through nucleation from an initial amorphous silicon layer after which the crystallites grow, the crystallinity is not constant throughout the layer if no attempt is made to control it, which may result in far from optimal layers. Therefore, research in microcrystalline Silicon deposition puts a lot of effort into crystallinity control during deposition through varying Si the deposition parameters. The bandgap of microcrystalline silicon depends on the fraction of amorphous silicon in the material. Layers with a substantial crystalline fraction have a bandgap close to that of crystalline Silicon (1.1 eV). The apparent higher absorption for such microcrystalline layers compared to single crystalline Si has been demonstrated to be caused by light scattering at the layer surfaces. The absorption below the bandgap is much higher than that for crystalline silicon, and is caused by defects within the bandgap. The absorption coefficient at those long wavelengths, therefore, gives a measure for the layer quality. It can be measured in different ways, but a powerful measurement technique that is increasingly be being used is the Fourier transform photocurrent spectroscopy (FTPS) [7]. Fig. 2-3 also gives important insight in layer quality degradation in microcrystalline silicon. When the first microcrystalline silicon solar cells were demonstrated, tests on devices with relatively high crystallinity led to the conclusion that microcrystalline Silicon did not suffer from light-induced degradation. As better devices closer to the transition were made and more detailed degradation studies were carried out, a more subtle picture has emerged: microcrystalline silicon suffers from a mild form of the photo-induced degradation [8]. The degraded values (after 1000 hours under standard illumination) in Fig. 2-3 show that, as expected, the degradation is worst for fully amorphous silicon, while it is negligible for almost fully crystalline layers. For microcrystalline layers in the transition region, there is a small yet significant degradation.

There is, however, no degradation at all if the high energy photons are filtered out [9], as is the case in tandem solar cells. μc-Si:H with a higher crystalline volume fraction than a-Si:H.

The assumption here is that the lower mobility band gap in μc-Si:H thin film cause the lower open circuit voltage. We would expect the free carrier densities to be higher in the bands for the “low Eg” μc-Si:H in which the band edges are closer to the quasi-Fermi levels. This is easily inferred from the expressions of the free carrier densities:

(E_C E _{F n} ) / K T

n = N C e ⁻ ⁻ Eq. 2-1

(E _{F p} E_V ) / K T

p = N V e ⁻ ⁻ Eq. 2-2

Fig. 2-3 Defect-related optical absorption in as-deposited and degraded state for a series of single-junction µc-Si:H silicon p-i-n solar cells as a function of the crystallinity of the intrinsic layer (i-layer) [8].

In the general, a μc-Si:H cell with a higher crystallinity has a higher Jsc, primarily because of higher free carrier mobilities and higher free carrier densities, which causes higher photo-generated hole trapping, therefore higher field near the P/I interface, and a collapse of the electric field over the volume and a lower Voc. In fact the above opinions are the reasons for a lower Voc and higher Jsc in μc-Si:H cells [11].

2.1.3 Tandem Solar Cell

Microcrystalline silicon has not only the advantage of a better stability under light exposure than amorphous silicon. Microcrystalline silicon has a different sunlight absorption spectrum than a-Si:H as shown in Fig. 2-4.

Fig. 2-4 The optical absorption coefficient α and the penetration depth dλ, where dλ = 1/α of monochromatic light with photon energy hυ and wavelength λ, for crystalline silicon (c-Si), and typical device-quality a-Si:H and µc-Si:H layers on glass. The curve for µc-Si:H has been corrected for light scattering due to surface roughness [10].

µc-Si:H absorbs the light more in the infrared range. This difference between a-Si:H and µc-Si:H has been used to increase the sunlight energy conversion efficiency by performing a so-called Tandem cell, which consists in a multi-junction cell with a top-cell made of a-Si:H and a bottom-cell of µc-Si:H. The thickness of the a-Si:H intrinsic layer is generally thin (about 250 nm) in order to collect a maximum of the photo-induced electrons. The µc-Si:H intrinsic layer has to be thicker (about 1.5 - 2 µm) because of its indirect band-gap and the necessity to match the photo-current generated by the two stacked cells [2].

From the point of view of the manufacturing technique, µc-Si:H is fully compatible with a-Si:H. Indeed, it can be deposited using the same reactor at the same substrate temperature of about 200 ^◦C. The source gases are silane and hydrogen as for the deposition of a-Si:H.

Generally, the hydrogen dilution is increased to change from the amorphous to the microcrystalline silicon deposition regime. When µc-Si:H is deposited in RF PECVD, the deposition rate (1 - 5 Å/s range) is generally less than for a-Si:H. This low deposition rate is a limiting factor for the use of µc-Si:H in PV solar cells, because it involves a too long time to deposit the ~ 2 µm thick µc-Si:H intrinsic layer. Typically, about 1.5 hours are necessary to deposit a 2 µm thick layer with a 4 Å/s deposition rate. In order to achieve low-cost Tandem PV solar cells, the deposition rate of the intrinsic µc-Si:H layer has to be improved to rates higher than 10 Å/s. High deposition rates are not the only condition for µc-Si:H films to achieve low-cost Tandem PV solar cells. Vetterl et al [12] have shown that the best material quality to be integrated in a Tandem solar cell is at the limit between amorphous and microcrystalline silicon. Therefore, the deposition conditions have to be perfectly controlled in order to have a material crystallinity in the 40 - 60 % range, and not to fall into the fully amorphous or strongly microcrystalline deposition region.

2.2 Deposition Mechanisms of µc-Si:H Thin Film

The mechanisms involved in the growth of hydrogenated microcrystalline silicon are complex, which combine both chemical and physical aspects. The chemical aspects include the hydrogen, silicon dangling bonds and reactions in the film. The physical aspects is such as the species (SiHx and H) bulk diffusion. The heterogeneous microstructure composed by crystalline grains embedded in an amorphous matrix adds to the system complexity. Moreover, the surface and sub-surface chemistry and physics depend on both the growing film properties and on the plasma. The three most important models are the surface diffusion model, the selective etching model and the chemical annealing model, which are reviewed in this section.

2.2.1 Surface Diffusion Model

The surface diffusion model was experimentally demonstrated as shown by Matsuda [46].

The crystalline volume fraction of deposited films strongly depends on the surface diffusivity of SiHx which is improved when the silane concentration is reduced or the substrate temperature increased. This is due to the hydrogen surface coverage improvement prolong the SiHx surface diffusion length to permit SiHx to attach at favorable sites to create a flat film surface. This has been confirmed by the abrupt fall of the film crystallinity for substrate temperature higher than 400 ◦C [13], caused by the desorption of hydrogen adsorbed on the surface at such high temperatures [14]. Moreover, some of the hydrogen atoms coming from the plasma recombine with surface bonded-hydrogen heating locally the surface, which enhances the surface diffusion of SiHx as sketched in Fig. 2-5. The role of the surface diffusion in flattening the film surface is of particular importance for the nucleation of

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