light trapping in thin film silicon solar cells on plastic substrates

Light Trapping in Thin Film SiliconSolar Cells on Plastic Substrates

Cover image: Microscope image of the grooves of ’Here comes the sun’ by theBeatles, on vinyl.Druk: Ipskamp Drukkers BV, Amsterdam

Light Trapping in Thin Film SiliconSolar Cells on Plastic Substrates

Lichtopsluiting in dunnelaagsiliciumzonnecellen op plastic substraten

(met een samenvatting in het Nederlands)

Proefschrift

ter verkrijging van de graad van doctor aan de Universiteit Utrechtop gezag van de rector magnificus, prof.dr. G.J. van der Zwaan,

ingevolge het besluit van het college voor promoties in hetopenbaar te verdedigen op woensdag 16 januari 2013

des ochtends te 10.30 uur

door

Micha Minne de Jong

geboren op 6 maart 1981 te Laren

Promotor: Prof.dr. R. E. I. SchroppCo-promotor: Dr. J. K. Rath

The work described in this thesis was financially supported by NL Agency(Agentschap NL) of the Ministry of Economic Affairs, Agriculture and Inno-vation of The Netherlands: program EOS-LT (Energie Onderzoek Subsidie -Lange Termijn).

Contents

1 Introduction 91.1 Renewable energy . . . . . . . . . . . . . . . . . . . . . . . . . . 91.2 Photovoltaic energy and solar cells . . . . . . . . . . . . . . . . 101.3 Silicon thin film solar cells . . . . . . . . . . . . . . . . . . . . . 121.4 Some basic solar cell physics . . . . . . . . . . . . . . . . . . . . 141.5 Low temperature flexible solar cells . . . . . . . . . . . . . . . . 161.6 Outline and objectives . . . . . . . . . . . . . . . . . . . . . . . 18

2 Experimental techniques 212.1 Silicon depositions: Plasma-enhanced chemical vapour deposition 21

2.1.1 The ASTER deposition system . . . . . . . . . . . . . . 222.1.2 The IRIS plasma characterisation system . . . . . . . . 24

2.2 Materials characterization . . . . . . . . . . . . . . . . . . . . . 252.2.1 Reflection-transmission measurements . . . . . . . . . . 252.2.2 Constant-photocurrent method . . . . . . . . . . . . . . 272.2.3 Raman spectroscopy . . . . . . . . . . . . . . . . . . . . 27

2.3 Solar cell characterization . . . . . . . . . . . . . . . . . . . . . 292.3.1 The solar simulator . . . . . . . . . . . . . . . . . . . . . 292.3.2 Spectral response . . . . . . . . . . . . . . . . . . . . . . 30

3 The role of temperature in plasma dust formation 313.1 Dusty plasmas: From α to γ’ . . . . . . . . . . . . . . . . . . . 313.2 The influence of temperature on dust formation . . . . . . . . . 333.3 Dust formation and OES . . . . . . . . . . . . . . . . . . . . . 34

3.3.1 Recording OES profiles . . . . . . . . . . . . . . . . . . 343.3.2 Dust formation as a function of power, hydrogen dilu-

tion, and temperature . . . . . . . . . . . . . . . . . . . 373.3.3 TEM images of dust . . . . . . . . . . . . . . . . . . . . 39

6 Contents

3.3.4 OES of pulsed Plasmas . . . . . . . . . . . . . . . . . . 413.4 Mass spectrometry . . . . . . . . . . . . . . . . . . . . . . . . . 42

3.4.1 Clusters, the precursors of dust formation . . . . . . . . 423.4.2 Ion energies . . . . . . . . . . . . . . . . . . . . . . . . 433.4.3 Cluster formation and temperature . . . . . . . . . . . 433.4.4 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . 47

4 Low temperature silicon layers 494.1 The role of substrate temperature in PECVD . . . . . . . . . . 494.2 Controlling the substrate temperature . . . . . . . . . . . . . . 50

4.2.1 Substrate stretch holder . . . . . . . . . . . . . . . . . . 514.2.2 Gas pressure . . . . . . . . . . . . . . . . . . . . . . . . 524.2.3 Plasma heating . . . . . . . . . . . . . . . . . . . . . . 54

4.3 Low temperature intrinsic layers . . . . . . . . . . . . . . . . . 564.3.1 a-Si:H intrinsic layers . . . . . . . . . . . . . . . . . . . 564.3.2 nc-Si:H intrinsic layers . . . . . . . . . . . . . . . . . . . 58

4.4 Low temperature doped layers . . . . . . . . . . . . . . . . . . . 614.5 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

5 Light trapping in amorphous silicon cells on polycarbonatesubstrates 655.1 Light trapping techniques . . . . . . . . . . . . . . . . . . . . . 65

5.1.1 Scattering . . . . . . . . . . . . . . . . . . . . . . . . . . 665.1.2 Nanopyramid periodic structures . . . . . . . . . . . . . 675.1.3 Geometric light trapping: micropyramid periodic struc-

tures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 685.2 Low temperature solar cells on PC substrates . . . . . . . . . . 73

5.2.1 Cells on PC: Experimental issues . . . . . . . . . . . . . 745.2.2 Solar cell results . . . . . . . . . . . . . . . . . . . . . . 77

5.3 Post-deposition treatments . . . . . . . . . . . . . . . . . . . . 835.3.1 Shunt busting . . . . . . . . . . . . . . . . . . . . . . . . 835.3.2 Post deposition annealing . . . . . . . . . . . . . . . . . 845.3.3 Stability under light soaking . . . . . . . . . . . . . . . 84

5.4 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85

6 Micromorph tandem cells on plastic substrates 876.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 876.2 nc-Si:H cells on glass substrates . . . . . . . . . . . . . . . . . . 886.3 Tandem cells on glass substrates . . . . . . . . . . . . . . . . . 91

6.3.1 Stability under light soaking . . . . . . . . . . . . . . . 93

Contents 7

6.4 Tandem cells on plastic substrates . . . . . . . . . . . . . . . . 956.5 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 99

Bibliography 101

Summary 115

Samenvatting 119

List of publications 123

Nawoord 125

Curriculum Vitae 127

Chapter 1

Introduction

1.1 Renewable energyThe worlds thirst for energy expands rapidly. The International Energy Agency(IEA) projects that the total energy consumption will grow 30% up to 2035compared to 2010, leading to a 20% rise in carbon dioxide (CO2) emission [1].90% of this increase in demand originates from emerging economies like Chinaand India, Brazil and the Middle East. These countries have a fast growingmiddle class population that consumes more and more energy and have a rapideconomic growth. The energy consumption and the growth of these countriesare dominated by fossil fuels. Of the energy consumed in 2010 world wide, 81%is generated from fossil fuel sources and this is estimated to be still 75% in2035. If this scenario becomes reality, the IAE estimates a global temperaturerise in excess of 3.5°C. Therefore, new non-fossil based energy sources need tobe explored.

A very important candidate for alternative energy is solar energy. Theamount of radiative power the earth receives from the sun is multiple thousandtimes the amount of energy consumed world wide. Therefore, if we are able toharvest this energy in an economically viable manner, solar energy could fulfila large fraction of the worldwide energy demand.

Solar energy generation can generally be divided in two different technolo-gies: Solar-thermal applications, in which solar radiation is used to heat up amedium, which in turn is used to power a turbine. In photovoltaics (PV), theradiation is directly converted into electrical energy.

If solar energy conversion is to become a large contributor to the energy

10 Chapter 1. Introduction

production, drastic cost reduction for solar cell modules is required. Solarenergy harvesting will only become an important means of energy productionif the price of solar energy can compete with conventional energy sources.As technology progresses, the production costs have decreased over the lastfew decades [2], while efficiencies have increased [3]. Also the upscaling ofproduction has been beneficial for cost reduction of solar panels. The price ofsolar energy not only depends on the price of the equipment, but depends alsoon the amount of sunlight available, which varies greatly at different locationson the globe. If solar energy is to compete with energy from fossil sources, alsothe price of other sources of energy plays a major role, which is also a locationdependent parameter. The point at which solar electricity can be harvestedat a price equal to that or lower than the price of electricity produced byconventional energy plants, is called grid-parity. If this point is reached, solarenergy becomes a viable alternative to conventional energy sources withoutsubsidy support, although government policies may still play a dominant role[4]. In some locations, solar energy harvesting has already reached grid parity[2].

1.2 Photovoltaic energy and solar cellsIn 1839, Alexandre-Edmond Becquerel discovered the photovoltaic effect [5].He observed that light could induce a current when an interface of two liquidswas illuminated. Nowadays, we are familiar with a range of semiconductorswhich we can use to convert photon energy into energetic charge carriers. Ifwe are able to extract these charge carriers from the semiconductor material,we can generate a current that can be used to drive an external circuit. Theseprinciples are the starting point of the development of solar cells.

When in 1954 Bell Laboratories reported on a p-n junction in silicon basedsolar cell with an efficiency of around 6%, great interest was born for thisconcept [6]. Initially, solar cells were far too expensive to be used for terrestrialelectricity generation, but became a standard source for space applications.Ever since, reported efficiencies have continuously been increasing whereas theproduction costs have decimated with production volume. In the last 10 years,PV is one of the fastest growing industries with annual growth rates more than40%. This increase is not only fueled by the progress in technology and lowermodule prices, but also by increasing prices for fossil fuel based energy, andto some extent by the awareness of the general public of the environment andclimate change and the government policies such as a feed-in-tariff for solarenergy [7].

1.2. Photovoltaic energy and solar cells 11

US$0.10/W US$0.20/W US$0.50/W

Thermodynamic limit

US$1.00/W

US$3.50/W

Single bandgap limit

Cost (US$/m2)

0 100 200 500 300 400 0

20

40

60

80

100

III

I II

Figure 1.1: Three generations of solar cells, showing their present and projec-ted cost and efficiency. From [9].

The largest fraction of the PV industries is crystalline silicon (c-Si) basedtechnology, with a market share of around 90%. Thin-film based technologiestake up about 10% of the market share [7, 8].

In PV technology, three generations are distinguished. First generationsolar cells are c-Si based single junction solar cells, which in terms of efficiency,will never cross the Shockley-Queisser (S-Q) limit of 30%. The S-Q limitdepends on the size of the band gap (Eg) of the used material [10] and is 34%for an optimum band gap of 1.4 eV. Crystalline silicon solar cells are based onsilicon wafers, which are sawed from single-crystal or multicrystalline siliconingots and have a typical thickness of a few hundred micrometers. Apartfrom the efficiency limit, a major drawback of this technology is the relativelyhigh material usage. Second generation solar cell technology is based on thinfilm technologies. The single junction type of 2nd generation solar cells haveagain suffered from S-Q efficiency limits, however, a multijunction cell can inprinciple lift the efficiency above the S-Q limit. Instead of bulk material asthe absorber material, these cells use thin films, which are deposited from thegas phase. This has a number of advantages. (1) Because the thin films have


a thickness in the order of micrometers or even thinner, only a fraction ofthe amount of material is needed. This makes them lighter, less fragile andsignificantly cheaper to produce. (2) The gas phase deposition processes allowfor large glass substrates [11, 12] to be used as substrate materials, copying thefabrication of displays. (3) Because other substrates than glass can be used,thin film solar cells can be made flexible, thereby making it possible to fabricatethem in a to a roll-to-roll manufacturing process, which may drastically cutproduction costs. Third generation solar cells can potentially break the S-Qlimit. Examples of third generation concepts are multi-junction cells [13], hot-carrier devices [14], spectral conversion techniques (up or down conversion)[15, 16] or quantum dot-based devices [17]. A schematic representation of thethree generations of solar cells is given in figure 1.1.

1.3 Silicon thin film solar cellsThe atoms in hydrogenated amorphous silicon (a-Si:H) do not form an orderedmatrix, unlike in c-Si, in which all atoms are fourfold coordinated. In a-Si:Hthere is a certain degree of variation in the bond length and bond angle distri-bution, which has implications for the band gap and leads to the presence ofband tail states. Because not all atoms are bonded to four neighbouring atoms,there will be non-bonding orbitals, the so-called dangling bonds. In a-Si:H ma-terial, most of these bonds will be passivated by bonding to incorporated Hatoms, although a number of dangling bonds will remain unpassivated, whichwill act as midgap states or electronic defects in the silicon matrix. In dopeda-Si:H layers, due to a thermal equilibrium between dopants and defect cre-ation, the defect density is even higher, leading to a higher recombination ratefor charge carriers. A conventional p-n junction, as used in c-Si solar cellswill therefore not perform adequately because most charge carriers are lostin the silicon layers through trapping and recombination before they can beextracted from the solar cell. Alternatively, by sandwiching an intrinsic a-Si:Hlayer between doped layers, a device can be created in which the intrinsic layeracts as the main light absorber. The p- and n-layers induce an electric field inwhich the charge carriers drift towards the contacts.

Crystalline silicon is an indirect semiconductor with a band gap of 1.1 eV.For every photon to be absorbed, momentum needs to be transferred to thesilicon lattice in the form of phonons to conserve momentum. Therefore theabsorption coefficient of c-Si is rather low and c-Si based cells are relativelythick. Due to the structural disorder in a-Si:H, the material acts as a directband gap semiconductor with an Eg of around 1.8 eV, which results in a higher

1.3. Silicon thin film solar cells 13

absorption coefficient. Therefore, a-Si:H solar cells can be made much thinnerthan their c-Si counterparts, with a typical i-layer thickness of smaller than500 nm.

In 1965, Sterling showed the possibility to deposit a-Si:H in a radio fre-quency (RF) discharge [18]. Chittick showed in 1969 that a-Si:H can have abroad range of photoconductivity, depending on the deposition temperature[19]. His group also showed the possibility of fabricating n-type a-Si:H by sub-stitutional doping by phosphorus. Ever since Spear discovered the possibilityof doping of a-Si:H both n-type and p-type by adding PH3 and B2H6 to the gasmixture in 1975 [20] and their following paper on the first a-Si:H p-n junctionin 1976 [21], numerous research groups and companies have been investigatingpossible applications of this material [22]. It was Carlson and Wronski [23]who made the first a-Si:H p-i-n solar cell, achieving a conversion efficiency of2.4%. With the introduction of the first multi-chamber system by Kuwanoet al. [24], which separates reaction chambers for the deposition of n-doped,p-doped and intrinsic layers to avoid dopant atom cross contamination, a newsolar cell record was set at 6.9% conversion efficiency. At present, the highestreported stabilized conversion efficiency of a single junction a-Si:H solar cellis 10.1% [3]. A major challenge for a-Si:H cells is the degradation under lightexposure, as first described by Staebler and Wronski in 1977 [25]. When thecells are exposed to light, the midgap defect density (Nd) increases, whichleads to degradation of the cell performance.

Another form of silicon that can be deposited from the gas phase is nano-crystalline silicon (nc-Si:H), which is also called microcrystalline silicon (µc-Si). Nanocrystalline silicon differs from amorphous silicon by its structure; itcontains grains of crystalline silicon, embedded in the amorphous tissue. Bychanging the atomic hydrogen concentration relative to silyl (SiH3) speciesreaching the growth surface, the crystalline fraction can be controlled. Nano-crystalline silicon shows properties different from amorphous silicon. Due tothe silicon crystallites, nc-Si:H has a lower band gap of ∼1.1 eV and thereforeshows a higher absorption in the red part of the spectrum, and can have ahigher electron mobility if there are not too many grain boundary defects.Furthermore, nc-Si:H solar cells are less sensitive to light induced degradation(LID). The material was first reported by Vepřek in 1968 [26] and uninten-tionally oxygenated nc-Si:H was first used to make a thin film solar cells in1992 [27], while in 1994 the first complete nc-Si:H cell was created by Meier etal. from IMT [28]. It was also the researchers from IMT who introduced the"micromorph" concept: a double junction solar cell, which comprises an a-Si:Htop cell and a nc-Si:H bottom cell to collect light from a broader spectrum[13]. At present, the highest reported and certified nc-Si:H single junction cell


has an efficiency of 10.1% [29]. The highest reported stabilized (light soaked)and certified efficiency for an a-Si:H/nc-Si:H tandem configuration is 12.2%[30]. Because a multi-junction cell is basically a set of solar cells connected inseries, the total current is limited by the cell that generates the lowest currentdensity. Therefore, the current generated by the cells needs to be matched,which imposes requirements on the thicknesses of the individual cells.

1.4 Some basic solar cell physics

Figure 1.2 shows a schematic representation of a superstrate type thin filmsolar cell in the p-i-n configuration. It generally consist of a substrate, a trans-parent front contact (transparent conductive oxide or TCO), a p-doped layer,an intrinsic (undoped) layer, an n-doped layer and a back reflector in succes-sion. Metal deposited contacts are used for carrier extraction from the frontand back side of the cell. The heart of the cell is made out of a semiconductormaterial, which can convert photon energy into excited charge carriers [5], ifthe energy of the photon is higher than the band gap of the semiconductor.In the case of a-Si:H, the material has an Eg of 1.8 eV at room temperature,with band-tail states, caused by weak Si–Si bonds, and mid-gap states, causedby silicon dangling bonds within the gap, caused by unbonded Si bonds in thematerial. When a photon is absorbed, this can lead to the formation of anelectron-hole pair, which can move through the semiconductor material. Thecharge carriers need to be moved to the external contacts by drift in an elec-tric field to be able to recombine in an external circuit, because a transportmechanism in the absorber layer based on diffusion alone would lead to a largerecombination loss. To accomplish this, an internal electric field is generatedby using doped layers. The absorber layer is made out of intrinsic silicon,because it has a midgap defect density nearly two orders of magnitude lower,compared to the doped layers.

Figure 1.3 shows the band diagram of a solar cell under short-circuit condi-tions (left) and under forward bias voltage conditions (right). Also shown arethe band gap, the quasi Fermi level (Ef ), and the electrons and holes. Apply-ing a forward bias (Vb) to the cell will reduce the electric field within the cell,causing a higher recombination rate. When there is no current, the cell is inopen-circuit conditions. The applied voltage at this point is the open-circuitvoltage (Voc).

1.4. Some basic solar cell physics 15

Substrate

TCO p-layer

i-layer

n-layer Back reflector Back contact

Front contact

Incoming light

Figure 1.2: Schematic presentation of the layers in a p-i-n superstrate typethin film solar cell.

Elec

tron

pot

entia

l

p i n p i n

Eg

Eg

+

-

+

-

Ef eVb

Figure 1.3: Band diagram of a solar cell under short-circuit conditions (left)and under forward bias conditions (right).


1.5 Low temperature flexible solar cellsFor solar energy generation to become a feasible alternative to fossil fuel-based electricity generation or to other means of renewable energy generation,the price of solar cells is extremely important. Thin film silicon solar cellshave the potential to be fabricated much cheaper than their crystalline siliconcounterparts. Not only the amount of material used is greatly reduced, also theenergy input to fabricate thin film solar cells is much smaller than for c-Si cells,resulting in a lower so-called energy payback time [31]. If a flexible materialis chosen for the substrate, bendable cells can be made, which have a numberof advantages over rigid substrates, such as glass, without loss in efficiency.First of all, plastics such as polyethylene terephthalate (PET), polyethylenenaphthalate (PEN) and polycarbonate (PC) are much cheaper than glass.More importantly, flexible substrates enable a roll-to-roll process, in which cellscan be produced in a continuous process, rather than a batch-type process,used for glass substrates. Especially when working with vacuum equipment,roll-to-roll processing can yield large cost reductions [32]. Furthermore, flexiblecells can be shaped in many ways, making them integrable into buildings andarchitecture. Using a lightweight and unbreakable substrate, such as plastics,savings can be made in transport costs. It has to be noted that the highestinitial efficiency of 16.3% for thin film silicon solar cells has been obtainedon a flexible substrate (stainless steel foil), using a triple junction (a-Si:H/a-SiGe:H/nc-Si:H) by Unisolar Ovonics [33].

The main challenge in depositing thin film solar cells directly on plasticflexible substrates are the limitations imposed on deposition temperature. Ifwe want to use cheap plastics such as PET, PEN or PC, we are limited to de-position temperatures of 70-100°C, 150°C or 130°C, respectively [34]. In 1974,Spear already observed changes in material quality of a-Si:H when changingthe substrate temperature in a plasma-enhanced chemical vapour deposition(PECVD) [35]. Figure 1.4 shows the defect density of a-Si:H and nc-Si:Hlayers as a function of substrate temperature. For both types of layers, aminimum in dangling bond density is observed at a substrate temperature ofaround 200-250°C [36], much higher than the temperatures permitted when us-ing cheap plastics as a substrate. Depositing at lower substrate temperatures,while keeping other deposition parameters constant, will produce layers with amuch higher defect density, resulting in lower solar cell efficiencies. Changingthe flow ratio of silane and hydrogen in the reaction chamber can be used tocompensate for these effects, but these measures will have direct consequencesfor the deposition rate or can change the phase of the grown material [37].Some of the changes in the material quality can be ascribed to the decrease

1.5. Low temperature flexible solar cells 17

Figure 1.4: Defect density in nc-Si:H and a-Si:H films as a function of substratetemperature. From [36].

of energy transferred to the growing film by ions in the plasma [38]. Keepingthe film thickness low will reduce the charge carrier recombination losses, inwhich case excellent light trapping schemes are needed for the solar cell.

Table 1.1 shows an overview of reported thin film silicon solar cells on flex-ible substrates. It is divided into two sections: cells in the p-i-n configurationand cells in the n-i-p configuration. The table shows that low deposition tem-peratures generally lead to lower conversion efficiencies. Cells fabricated in atransfer process are deposited on a temporary substrate that is able to with-stand high temperatures and then transferred to a plastic substrate, enablinghigh deposition temperatures. Most efforts at depositing low temperature cellshave focussed on n-i-p substrate type cells. Because n-i-p cells do not requirea transparent substrate, a wider range of substrate materials can be used.

To fabricate low-cost thin film solar cells, not only the material costs ofthe TCO, absorber material and the substrate material are important, alsothroughput is important. For micromorph tandem cells, the bottom nc-Si:Hcell has a typical thickness of 1 to 3µm, which has an impact on deposition


Superstrate (p–i–n) cells :Cell type D/T Substrate T (°C) η (%) Sourcea-Si:H D PET 110 4.9 [39]a-Si:H T Polyester 200 7.7 [40]

a-Si:H/nc-Si:H T Polyester 200 9.1 [41]Substrate (n–i–p) cells :

Cell type D/T Substrate T (°C) η (%) Sourcea-Si:H/a-SiGe:H/nc-Si:H D SS n/a 16.3 [33]

a-Si:H/a-SiGe:H D Kapton n/a 10.1 [42]a-Si:H D PEN 150 8.7 [43]a-Si:H D E/TD 140 6.0 [44]nc-Si:H D LCP 180 8.1 [14]a-Si:H D PET 100 5.9 [45]

E/TD: ethylene–tetracyclododecene copolymer, SS: stainless steel, LCP: liquid crystal polymer

Table 1.1: Present reported record initial efficiencies of silicon thin film solarcells deposited on flexible substrates for p-i-n and n-i-p types of cells. D or Tdenotes a direct deposition or a transfer-process.

times and therefore on manufacturing costs. Moreover, the top a-Si:H cellhas to be thin, because thick a-Si:H cells are more sensitive to light induceddegradation than very thin a-Si:H layers. For these reasons, thin absorberlayers should be used. However, to ensure good light absorption within theactive layers, light trapping schemes must be deployed. These schemes havethe purpose to enlarge the optical path of travelling photons in the cell, whilekeeping the electrical paths short, enabling a good light absorption withoutsacrificing too much on electrical performance. The traditional way in thinfilm (superstrate) p-i-n solar cells is to create a rough interface between theTCO and the p-layer of the cell, by texture etching the doped ZnO in an acidicsolution or depositing natively textured SnO2 [46, 47, 45]. Other schemes suchas rough silver, 1D or 2D [48] gratings or 3D nanopillar designs [49, 50] haveshown to improve the generated current in n-i-p cells.

1.6 Outline and objectivesThe goal of this thesis is to investigate and find solutions for the difficulties thatare encountered when depositing thin film silicon solar cells on plastic flexiblesubstrates. On the level of deposition plasmas, we investigated the influence

1.6. Outline and objectives 19

of the deposition temperature on a number of plasma characteristics, with afocus on the formation of dust, which is a temperature dependent process. Theinfluence of deposition temperature on the quality of silicon layers is studied.It was investigated how the deleterious effects of low deposition temperatureon material quality can be compensated. Finally, solar cells were made at lowtemperatures, both on glass substrates and on plastic substrates. Differentlight trapping management techniques were tested for amorphous silicon solarcells, nanocrystalline solar cells, and micromorph tandem solar cells.

Chapter two introduces the main experimental techniques that are used forsilicon and metal oxide depositions, optical and electrical material character-ization and solar cell characterization techniques which are used for the studyfor this thesis.

Chapter three is concerned with the changes that are induced in the plas-mas when changing the deposition temperature. Using a newly developedtechnique that utilizes the axial optical emission profile from the plasma wewere able to identify whether a plasma is dust-free or produces dust particles.Furthermore, the influence of the substrate temperature is investigated on thefirst stage of dust formation: cluster formation. Using mass spectroscopy, theformation of polysilanes is studied as a function of substrate temperature.

The subject of chapter four is the study of the optical and electrical prop-erties of amorphous and nanocrystalline silicon deposited at substrate temper-atures below the optimum temperature and to develop plasma conditions atwhich device-quality silicon layers can be deposited.

Chapter five treats the deposition and characterization of amorphous sil-icon thin film solar cells deposited at low temperature (130°C) on glass and thedeposition of these low temperature cells on plastic polycarbonate substrates.First, a number of practical concerns are treated: the adhesion of layers to theplastic substrates, the curving of plastic substrates due to thermal expansionand the degassing of plastic substrates. A number of light trapping schemeswas studied: Scattering by a rough interface between the front contact (TCO)layer and the p-layer, geometric light trapping by pyramids that are largerthen the effective wavelength of light in the material and finally light manage-ment by pyramidal structures comparable to the effective wavelength of light.It was investigated how these texturization schemes enhance the light absorp-tion, but also their influence on the electrical quality of the cells: Voc and fillfactor. Furthermore, a number of post-deposition treatments are investigated:thermal annealing, shunt busting and light induced degradation.

The last chapter presents nanocrystalline solar cells and amorphous/ nano-crystalline silicon (micromorph) tandem solar cells deposited at low substratetemperature (130°C) both on glass and on polycarbonate substrates. First


nanocrystalline solar cells are deposited on glass and characterized. a-Si:H/nc-Si:H tandem cell are fabricated at low temperature on glass and on plasticsubstrates.

A part of this work was done in collaboration with Wageningen Univer-sity Glastuinbouw, who perform research on the use of micro- and nano tex-tured glass and plastics for ultra-transparant greenhouse roofing to increasecrop production, and with Aquamarijn Micro Filtration BV, who provided themicro-pyramid structured polycarbonate substrates.

Chapter 2

Experimental techniques

This chapter describes the experimental techniques that were used to depositdifferent layers and solar cells used in this thesis and the techniques to char-acterize them.

2.1 Silicon depositions: Plasma-enhanced chem-ical vapour deposition

Plasma-Enhanced Chemical Vapour Deposition (PE-CVD) is a form of CVDthat can be used for thin film depositions. As opposed to most other forms ofCVD, PE-CVD can be operated at low temperatures. In the process, a sourcegas is dissociated in an oscillating electric field between two parallel plates. Inour case, the substrate is mounted to the grounded electrode. Between theplates, a source gas is introduced. For silicon thin film depositions, usuallySiH4 and H2 are used. In the electric field, electrons are accelerated and maycollide with a gas molecule. If this collision is sufficiently energetic, the impactcan cause ionization of the molecule, thereby creating an extra free electron,which can in turn collide with a molecule. This avalanche of reactions willresult in a plasma containing (positive) ions and free electrons. Because theelectrons are much lighter than the ions, the electrons are faster and will becollected at the electrodes, resulting in a positive plasma bulk.

Traditionally an excitation frequency of 13.56 MHz is used, due to legalrestrictions on the use of other radio frequency bands. Changing the excitationfrequency changes the plasma properties, such as the ion energies and biasvoltage in the plasma [51], which in turn can be beneficial or detrimental for

22 Chapter 2. Experimental techniques

the layer quality of the grown silicon layers.

2.1.1 The ASTER deposition system

The silicon layer depositions in this thesis are performed in the ASTER (Amorph-ous Semiconductor Thin-film Experimental Reactor) ultra-high vacuum mul-tichamber deposition system [52], using Very High Frequency PE-CVD (VHFPE-CVD). The chamber features 5 deposition chambers, mounted to a centraltransport chamber: one for p-type silicon depositions, one for n-type depos-itions, two for intrinsic silicon depositions and one experimental reactor fornanocrystal formation and deposition. A parking chamber is used for stor-ing under vacuum and gradual cooling of samples. Samples can be up to10× 10 cm2 and are mounted on a titanium substrate holder. The holder isinserted into the system through a loadlock chamber and can be transpor-ted between the separate chambers by a robot arm in the central chamber.The deposition chambers are equipped with viewports to monitor the plasma,either visually or using a spectrometer. In the intrinsic silicon chambers theinter-electrode distance can be changed from 5 to 27mm. These reactors areequipped with showerhead-type powered electrodes for an even distribution ofthe feedstock gasses into the plasma. All intrinsic layers are deposited at aVHF excitation frequency of 60Mhz, whereas the doped layers are grown at50MHz. The impedance of the plasma reactor can be matched to the 50Ω ofthe power input system through a set of adjustable capacitors, which form aan L-type matching network. The amount of reflected power (measured with aRhode&Schwarz NAP power meter) can be minimized to less than 1% of theinput power. The area of the powered electrode is 170 cm2 in the chambers forintrsinsic material deposition 150 cm2 in the chambers for doped depositions.

As source gasses for the depositions we use silane (SiH4) and molecularhydrogen (H2) for the growth of intrinsic silicon layers. By changing the ratioof the two source gasses, we can control the phase of the silicon to be eitheramorphous, nanocrystalline or mixed-phase. For doped layers, dopant gassesare added to the gas mixture. Trimethylboron (B(CH3)3 or TMB) is addedfor p-type doping and phosphine (PH3) is used for n-type doping.

Optical Emission Spectroscopy

During the deposition the process can be monitored in situ by optical emissionspectroscopy (OES). If a molecule A (or AB) is excited by electron impacteither by direct excitation

2.1. Silicon depositions: Plasma-enhanced chemical vapour deposition 23

2 0 0 3 0 0 4 0 0 5 0 0 6 0 0 7 0 0 8 0 0 9 0 00

1

2

3

4

Hα

S i H * Hβ

Inten

sity (

a.u.)

W a v e l e n g t h ( n m )

S i *

Figure 2.1: An example OES spectrum recorded from one of the ASTERdeposition chambers of a silane-hydrogen plasma. The different line emissionsare identified.


A+ e− → A∗ + e− (2.1)

or by dissociative excitation

AB + e− → A∗ +B + e− (2.2)

the excited molecule A∗ can relax to its ground state

A∗ → A+ hν (2.3)

emitting a photon. By recording the wavelength of the photons, we can identifythe species in the plasma. Species of interest are Si∗(at 289 nm), SiH∗ (at414 nm), Balmer alpha (Hα, 656 nm) and Balmer beta (Hβ , 490 nm). Theemission rates are associated with the dissociation rate of the different spe-cies. An example OES spectrum, recorded in ASTER from a silane/hydrogenplasma, is presented in figure 2.1, showing the emitted lines. Furthermore,information about the electron temperature can be extracted from the OESinformation.

The light, emitted by the plasma, is monitored through a viewport andan optical fibre and analysed and recorded using an Avantes MC2000 spectro-meter. The peaks found in the emission spectrum are fitted to Gaussians aftersubtracting a local background.

To reduce deposition on the viewport window, which would influence thetransmission of the window, it is shielded from the plasma using a valve whenno measurements are taken. An assembly of two horizontal slits is used toobtain a vertical emission profile of the plasma. This technique is used forthe detection of dust in the plasma, as described in chapter 3. A schematicrepresentation of one of the ASTER deposition chambers and the OES setupis given in figure 2.2.

2.1.2 The IRIS plasma characterisation systemThe IRIS (Ions and Radicals in Silane plasmas) system is designed to examineions and molecules formed in the plasma. Therefore it is equipped with aHiden EQP300 mass spectrometer. The plasma chamber is a copy of one of thedeposition chambers from ASTER, but at the position of the substrate (at thegrounded electrode), an orifice is fitted that leads to a separate chamber, whichis pumped to an ultra-high vacuum. Behind the orifice, the mass spectrometeris mounted. Ions and radicals that would normally reach the growing surfaceat the substrate will now travel through the orifice, into the mass spectrometer.

2.2. Materials characterization 25

VHF Gas inlet

Reactor wall

Heater/ substrate holder

Optical fibre

Optical port

Optical slits

Showerhead electrode

Shutter Movable platform

Quartz window

VHF Plasma

Figure 2.2: A schematic representation of one of the ASTER deposition cham-bers and the OES setup.

A set of electrostatic lenses leads the ions into the mass spectrometer to obtainoptimal yield. The mass spectrometer is not only mass-sensitive, but can alsodistinguish between species of different energies. The chamber behind theorifice is differentially pumped to make sure that the species do not collidebehind the orifice, thereby changing their energy and trajectory.

If neutral species are to be detected, they need to be ionized before theyenter the mass spectrometer. This is achieved by ionizing them by an electronemission filament. The system is used for the detection of dust precursors, asdescribed in chapter 3 and described in more detail by E. Hamers [53]. Thesystem can also be fitted with an OES system, similar to the system attachedto the ASTER system.

2.2 Materials characterization

2.2.1 Reflection-transmission measurementsThe R-T mini setup

The Reflection/Transmission setup (RT) from M. Theiss hard- and Software[54] measures the specular reflection and transmission simultaneously on thesame spot on a sample. A halogen lamp, sample stage, connected with anoptical fibre to a spectrometer enable recording of the spectra from 380 to1050 nm. The data can be analysed using a software package called ’SCOUT’


VHF

Gas inlet

Reactor wall

Heater

VHF electrode

Orifice

To Hiden mass spectrometer

Differentially pumped chamber

Reaction chamber

VHF Plasma

Figure 2.3: A schematic representation of the IRIS reactor chamber and theHiden mass spectrometer attachment.

by W. Theiss [54], which can use several models for the calculation of thewavelength-dependent absorption coefficient (α) and refractive index (n) fordifferent materials. These properties are then used to simulate transmissionand reflection spectra for layer stacks of different materials and different thick-nesses. These can be fitted to the measured spectra. In the case of silicon, weuse the O’Leary, Johnson and Kim model (OJL) [55]. This model describesmathematically the shape of the valence band and conduction band densitiesof states of semiconductors, which can be used to calculate the absorption coef-ficients. The Kramers-Kronig relation is used to calculate the refractive index(n) . By fitting these models to the acquired data we obtain the wavelengthdependent α and n of the measured material. From this information we candistil the optical band gap of the material.

The UV-VIS spectrometer

Whereas the R-T mini setup only measures specular transmission and reflec-tion on a sample, the Perkin Elmer Lambda 2S UV-VIS spectrometer setup isequipped with an integrating sphere, which can also measure diffuse reflectionand transmission, which can, for instance, be important for light scatteringproperties of substrates. In this way we can discriminate between specular,diffuse and total reflection and transmission. Because the wavelength range of

2.2. Materials characterization 27

this apparatus is between 200 and 1100 nm, we can measure in the ultraviolet(UV) range, in constrast to the R-T mini setup. Whereas the R-T mini setupmeasures reflection and transmission on the same position on the sample sim-ultaneously, this is not possible in the Perkin Elmer, because it needs to besetup differently for transmission and reflection measurements.

2.2.2 Constant-photocurrent methodThe sub-band gap absorption of silicon holds information on the slope of theband edges of the band diagram and on the density of mid-gap defect states.A method to measure the sub-band gap absorption is the use of the constant-photocurrent method (CPM). This technique is based on the photoconductiv-ity of the sample when it is illuminated at a certain wavelength. The light issupplied by a 250W halogen lamp and guided through a filter wheel, contain-ing 25 different interference filters in the red and infrared part of the spectrum.For every filter, the lamp intensity is changed until the photocurrent matchesa predefined value, while (a relative measure for) the number of photons (Nph)reaching the sample is recorded. Now the optical absorption coefficient α isproportional to 1/Nph. The absolute absorption coefficients can be found bycalibrating the CPM absorption coefficient to to the absorption coefficientsfound in RT measurements for λ< 1000 nm. For a-Si, from the absorptioncoefficient at hν=1.2 eV (α1.2) the defect density Nd can be calculated fromNd=Fα1.2, where F is a calibration factor, found to be 1016cm−2 [56]. Theband-edge absorption shows exponential behaviour, called the Urbach tail,which can be expressed as

α(λ) = α0e−E(λ)/E0 (2.4)

where α0 is a proportionality factor, E(λ) is the photon energy and E0 is theso-called Urbach energy, which is obtained by fitting a logarithmic slope tothe absorption coefficient as a function of photon energy.

2.2.3 Raman spectroscopyThe crystalline volume fraction of a sample can be quantified by Raman spec-troscopy measurements. In this measurement, the sample is locally illumin-ated by a strong laser. Besides the light that will be absorbed, transmittedor specularly reflected, a portion of the light will be inelastically scattered. Asmall fraction of the light shows a frequency shift (the Stokes shift) causedby interactions with phonons in the material. This phenomenon is called


250 300 350 400 450 500 5500

1000

2000

3000

4000

5000

6000

Raman shift Hcm-1L

Inte

nsity

Hcoun

tsL

Figure 2.4: An example of a Raman spectrum (dots) and the different fittedGaussians (solid lines) of a mixed phase thin film silicon layer deposited on aPC substrate.

Raman scattering and gives information about the density of states of thedifferent phonons. The different measured modes in the case of silicon arethe transverse-optic mode (TO, 520 cm−1), associated with crystalline siliconand the transverse-acoustic (TA, 100-200 cm−1), longitudinal-acoustic (LA,300-360 cm−1), longitudinal-optic (LO, 380-450 cm−1) and the transverse op-tic (470-490 cm−1) modes associated with amorphous silicon [57, 58].

In our setup, a Spectra Physics Ar+-ion laser (514 nm) illuminates thesample at an angle of 30°. The light is polarized horizontally before hittingthe sample. After hitting the sample, backscattered light is focussed througha set of lenses into a Spex triple monochromator and recorded by a nitrogencooled Roper Scientific CCD camera, after being polarized vertically. Thedata is analysed by fitting a number of Gaussian peaks to the spectrum: 3peaks fixed at 330, 440 and 480 cm−1 for the amorphous phase and two peaksbetween 505 and 520 cm−1 for the crystalline phase. The area of these peaksis then divided, resulting in the Raman crystalline fraction [59]:

Rc = I510 + I520

I480 + I510 + I520(2.5)

where Ix denotes the integrated intensity of the fitted peak at x cm−1. Al-though the Rc gives a quantitative measure for the crystalline fraction, it doesnot correspond exactly to a volume fraction. Nevertheless it is a useful quant-ity to compare the crystalline fraction of different samples. Figure 2.4 shows

2.3. Solar cell characterization 29

an example of a Raman spectrum, after background subtraction, measuredon a mixed phase a-Si/nc-Si silicon thin film deposited on a PC substrate atlow temperature. Also shown are the different Gaussians fit to the spectrum,associated with the different phonon modes.

2.3 Solar cell characterization2.3.1 The solar simulatorTo test the conversion efficiency (η) of solar cells, we can perform current-voltage (IV) measurements under standard illumination conditions. Interna-tionally it is agreed that these measurements are performed under 100mW/cm2,AM1.5 illumination [60] at 25°C. In our setup the light is produced by a com-bination of a xenon and a halogen lamp, through a set of mirrors and lenses.In the ideal case, the solar cell shows the behaviour of a diode, in parallel witha current source Jph and a resistance Rp and in series with a resistance Rs,which can be described by

J(V ) = −Jph+J0(exp(e(V − JRs)ndkT

)− 1) + V − JRsRp

(2.6)

where J0 is the reverse saturation current, e is the electron charge, nd is thediode quality factor, k is Boltzmann’s constant and T is the temperature inKelvin. From the measurements under illumination we can obtain the shortcircuit current density (Jsc), which is the current when V = 0, the open circuitvoltage (Voc), which is the voltage when J=0 and the fill factor (FF), whichis defined as

FF = JmppVmppJscVoc

(2.7)

where Jmpp and Vmpp are the current and voltage at the point where theproduct of the current and voltage peaks, the maximum power point. Thenthe conversion efficiency is

η = PmppPlight

= JmppVmppPlight

= FF JscVocPlight

(2.8)

in which Plight is the incident light power density, which in the case of AM1.5light is 100mW/cm2. The slope of the curve at V=0 is associated with Rp,whereas the slope at J = 0 is associated with Rs.

When IV characteristics are performed under dark conditions (Jph = 0) wecan extract the diode quality factor nd and the reverse saturation current J0.


2.3.2 Spectral responseThe Spectral Response (SR) of a solar cell tells us the fraction of photons thatis converted to an electron that reaches an external circuit as a function ofphoton wavelength, also called the quantum efficiency or external collectionefficiency. To measure the SR, modulated light of a xenon lamp is led througha monochromator and led onto the sample, for which the generated currentis measured using a lock-in amplifier. The light incident on the sample iscalibrated using a photodiode with a known spectral response. The SR can bemeasured from 350 up to 1100 nm and is equal to

SR(λ) = Iph(λ)nφ(λ) e (2.9)

where Iph is the measured photocurrent, nφ is wavelength dependent amountof photons directed at the sample per second and e is the electron charge.

From the wavelength dependent SR, the Jsc under AM1.5 illumination canbe calculated:

Jsc = e

ˆSR(λ)φAM1.5(λ) dλ (2.10)

in which φAM1.5(λ) is the wavelength dependent photon flux.Commonly, an externally applied bias voltage is used during measurements.

A negative bias voltage is used to strengthen the internal electric field of thesolar cell to reduce the carrier recombination in the i-layer and thus measurethe maximum current generating capabilities of the cell. A positive bias voltagecan be used to investigate the cell’s performance under maximum power point-conditions. Bias light that resembles AM1.5 light (both in spectrum and in-tensity) can be applied to obtain SR under standardized operation conditions.

Because a tandem cell consists of two cells connected in series, the max-imum measured current at a certain wavelength is limited by the cell thatgenerates the least current at that wavelength. Therefore, the individual cellswithin a tandem solar cells can be measured by illuminating the cell with con-tinuous (i.e. not chopped) bias light of different colours, thereby ’activating’the cell that responds to that colour, such that we can measure its counterpartby making sure that is does not limit the current from the cell that is beingtested with chopped monochromatic illumination. For this purpose, the biaslight can be led through a number of different filters: a red filter for measuringthe top cell and a blue filter for measuring the bottom cell in a tandem cellconsisting of an a-Si top cell and a nc-Si bottom cell.

Chapter 3

The role of temperature inplasma dust formation

3.1 Dusty plasmas: From α to γ’A challenging problem in low temperature depositions of thin-film silicon lay-ers is the formation of dust particles in the plasma. When these particles areincorporated in the silicon layers, they can introduce voids which will increasethe disorder in the amorphous network and will thus introduce electronic de-fects. Because the particles can be large compared to the film thickness, thedust particles can also cause electrical shunts through the layers [61], althoughrecent research shows that controlled dust formation in the plasma can bebeneficial for device performance when the layers are deposited at high rate[62].

The process of dust formation can be divided into three phases: startingwith a dust free plasma (the so-called α-regime), in the first phase, negativelycharged clusters may form through polymerization reactions in the plasma.Figure 3.1 shows a schematic representation of the potential profile in thebulk of the plasma. The bulk has a positive time-averaged potential, whichcauses negatively charged clusters to be trapped inside the plasma bulk. If theclusters collide with positively charged ions, they will become neutrals and willleave the plasma by diffusion, unless they collide with electrons before theyleave the plasma, thereby collect negative charge and remain trapped. Becausethe electron-capture cross section depends strongly on the size of the clusters[63], only clusters that are large enough (> 2nm) will be trapped inside the

32 Chapter 3. The role of temperature in plasma dust formation

Figure 3.1: Schematic presentation of a potential profile in a VHF plasmareactor through a VHF cycle. The dashed and dotted lines show the plasmapotential at φ=0.5π and φ=1.5π, respectively. The solid line shows the timeaveraged potential. In this graph, the grounded electrode is situated on theleft whereas the powered electrode is on the right.

3.2. The influence of temperature on dust formation 33

plasma bulk. In the second phase, when these clusters reach a critical size andconcentration, the clusters will coagulate and dust will start to form [64, 65].After coagulation has taken place, the particles have a typical size of a few tensof nanometres and will quickly acquire negative charge and therefore repel eachother, preventing further aggregation. The last phase represents the growing ofthe individual dust particles by the attachment of neutral particles or positiveions. The dust particles are now either lost at the sides of the reactor, due toa net gas flow towards the pumps or by diffusion, or are trapped in the reactoruntil the plasma is switched off. When dust has formed that remains trappeduntil the plasma is switched off, the plasma is in the so-called γ′-regime.

3.2 The influence of temperature on dust form-ation

The substrate temperature has a direct influence on the gas temperature insidethe reactor. It has been shown that in our system, in the temperature rangeused, the average (including the inactive parts of the reactor) gas temperaturerises around 25° when the substrate temperature is increased 100° [66]. Ithas been reported that choosing a higher substrate temperature can suppressdust formation. Although the lower gas density at higher gas temperaturemay play a role, the main mechanism is believed to be the dependence of thepolymerization rate of negative clusters on the gas temperature [63, 67]. Whenthe gas temperature decreases due to a decrease in the substrate temperature,the critical cluster size and concentration for coagulation can be reached muchquicker. Therefore, especially at low substrate temperatures, it is importantto monitor the dust formation in the plasma during the deposition process andto identify parameter windows for the deposition of dust free silicon. Laserlight scattering experiments, which measure the reflected light from dust inthe plasma, have shown to be a powerful tool to study the last phases ofthe dust formation process [68, 69]. Recent investigations show nanoparticlecharacterization using white light [70]. However, one of the drawbacks of theseoptical techniques is that the deposited silicon films on the viewports absorba part of the light, resulting in a time-dependent signal [68]. A method thatuses no optical detection is the spectral analysis of the radio-frequency currentto monitor dust formation. By measuring the amplitude of the fundamentaland the higher harmonics of the current through the plasma the productionof nanometre sized particles can be detected [71, 72].

In this chapter we present a non invasive in-situ diagnostic tool for mon-


itoring dust formation, based on optical emission spectroscopy (OES). By re-cording the optical emission lines for several species in the plasma as a functionof the vertical position between the electrodes we construct the emission pro-file of the optical emission of the plasma. Using these profiles we are ableto identify the plasma regime. The advantage of this technique is that thestate of the plasma (dust free or dusty) is marked by the asymmetry of theOES intensity profile and not by the absolute intensity value. Therefore thistechnique is insensitive to the loss of transmittance of the viewport due tosilicon film deposition, which can be a disadvantage when using other opticaltechniques. Because optical emission can also be used to predict the materialphase [73], this study shows that a single technique can be used as an in-situ plasma diagnosis tool for monitoring of the amorphous to nanocrystallinetransition as well as the transition of the dust-free to the dusty regime withouta supplement technique. Therefore, it is a plasma monitoring tool to controlthe complete silicon processing of “micromorph cell” manufacturing.

3.3 Dust formation and OES

3.3.1 Recording OES profilesThe presence of dust has an influence on the optical emission from a plasma.After the coagulation phase, when particles in the plasma can have sizes of tensof nanometres, their cross-section for capturing electrons is greatly increased.This has a direct effect on the plasma properties. The electron density dropsan order of magnitude [63], while the electron temperature is greatly increased.These properties are reflected in the optical emission from the plasma, as moreelectrons have sufficient energy to excite the different plasma species. Becausethe electron density in a plasma discharge can be significantly non-uniform[74], space-resolved OES measurements can provide information on the localelectron temperature of the plasma. If we compare the optical emission froma dust-producing plasma to a dust-free plasma, the local changes in emissionwill tell us where the dust is located. Whereas for atomic species or smallclusters gravity can be neglected when compared to other forces like elec-trostatic forces or thermophoresis, larger dust particles will be influenced bygravity and therefore be pulled towards the sheath at the bottom poweredelectrode.

To record axial emission profiles from the plasma, we used an AvantesMC2000 spectrometer connected to an optical fibre positioned behind an as-sembly of two horizontal slits of 1mm wide placed at a distance of 80mm

3.3. Dust formation and OES 35

0 5 1 0 1 5 2 0 2 5 3 002468

1 01 2

Optic

al Em

ission

(a.u.

)

V e r t i c a l p o s i t i o n ( m m )

S i H * H

β 3 x

Hα 2 x

P o w e r e d e l e c t r o d e

G r o u n d e d e l e c t r o d e

Figure 3.2: Cross-sectional optical emission profile of a plasma in the α-regime.Shown are the SiH∗, Hα and Hβ lines. The position is measured from the uppergrounded electrode downwards.

from each other and 20 cm from the plasma centre. A quartz window is usedto ensure transmission in the ultra-violet. This system is mounted on a stagethat can be moved in the vertical direction. The position is measured from theupper grounded electrode downwards. The spectral range of the spectrometeris 200 nm to 900 nm. A schematic representation of the setup is given in figure2.2. Using this system we recorded horizontal slices of the optical emissionfrom the plasma, with a spatial resolution of 1mm. From the recorded spectrawe derived the relative intensity of the lines associated with different plasmaspecies, by subtracting a local background and fitting Gaussians to the peaksfound in the spectra. In this way we obtained the relative signal intensity ofBalmer-alpha (Hα), Balmer-beta (Hβ), excited SiH (SiH*) and excited Si (Si*)as a function of vertical position.

Because it can take several minutes before the α to γ′ transition takes place,we waited for the emission to become stable before recording the spectra.

Figure 3.2 shows a typical optical emission profile for various emission linesof a plasma in the α-regime in the ASTER deposition system. The plasmaparameters were P = 13W (powered electrode area 170 cm2), R =45 andTs = 200°C, where P is the applied power, R is the hydrogen flow dilution


0 5 1 0 1 5 2 0 2 5 3 002468

1 01 21 41 61 8

P o w e r e d e l e c t r o d e

Optic

al Em

ission

(a.u.

)


S i H * H

β 3 x

Hα 2 x

G r o u n d e d e l e c t r o d e

Figure 3.3: Vertical optical emission profile of a plasma in the γ’-regime.Shown are the SiH∗, Hα and Hβ lines. The position is measured from theupper grounded electrode downwards.

ΦH2/ΦSiH4 and Ts is the substrate temperature. The inter-electrode distancewas 27mm. This axial emission distribution is very typical for particle-freeplasmas: Maxima in emission in the plasma sheath near both electrodes anda minimum in emission in the centre [74, 75]. The somewhat higher emissionnear the bottom powered electrode can be ascribed to the asymmetric designof our reactor where the powered (bottom) electrode has a smaller area thanthe grounded (upper electrode+chamber wall) electrode. Figure 3.3 shows theoptical emission profile for a plasma using the same deposition parametersbut at an applied power of 16W. This plasma is in the (dusty) γ’-regime,which is confirmed (not shown) by a shift in the impedance towards a moreresistive plasma [76]. The profile changes from two rather symmetric peaks atthe plasma sheaths and low bulk emission to a large peak in emission at thesheath near the lower (powered) electrode and a smaller peak at the upper(grounded) electrode, along with higher bulk emission. Pulled by gravity,the dust particles accumulate near the bottom electrode, where gravity iscounteracted by the force that the negatively charged particles experience fromthe potential drop near the electrode. The increase in optical emission in thislower region of the plasma in the γ’-regime can be ascribed to the presence of


dust particles. Because the dust particles act as electron traps, the electrondensity decreases and therefore the energy per electron increases, which inturn enhances the emission. The emission intensity is proportional to the rateconstant of emission by the relation ISi=KSiH4 Ne.NSiH4 , where Ne and NSiH4

are the electron concentration and silane concentration respectively and KSiH4

is the rate constant that depends on the electron temperature. The increase ofelectron temperature in the γ’ regime therefore increases the optical emissionintensity, especially at the powered electrode.

For large monodisperse injected particles in a non-reactive plasma, thiseffect is limited mainly to the region close to the bottom electrode [74]. Becausein our reactor the dust is grown rather than injected, we expect a variety insize and mass and therefore a broader axial distribution of the dust. Some ofthe lighter dust particles will be located near or in the plasma bulk, therebyalso enhancing the optical emission from the bulk of the plasma.

Because we can assume a Maxwell-Boltzmann distribution for the electronenergy for a low-pressure plasma [77] and because the electron temperature iswell below the minimum electron energy to excite hydrogen for emission in theBalmer series, the Hα and Hβ emissions can be ascribed to excitation of hy-drogen by electrons in the high energy tail of the electron energy distribution.Because of the differences in excitation energies (16.0 eV for Hα and 16.6 eVfor Hβ [78]), the ratio of the intensities Hβ/Hα can be used as a qualitativemeasure for the electron temperature [79].

Figure 3.4 shows the spatially resolved Hβ/Hα emission ratio for a plasmain the α-regime and for a plasma in the γ’-regime at the two above men-tioned deposition conditions. We observe an increase in electron temperaturethroughout the plasma reactor in the γ′-plasma, which is most pronouncedin the bulk of the plasma. This again reveals the presence of trapped dustparticles in the plasma.

3.3.2 Dust formation as a function of power, hydrogendilution, and temperature

Whether dust is produced in a plasma depends on the conditions under whichthe plasma is maintained. It has been shown before that increasing power dens-ity input, decreasing hydrogen dilution, increasing gas pressure or lowering thesubstrate temperature can change a dust free plasma into a dust producingplasma [80, 81, 82]. Using our optical method we have investigated the influ-ence of applied VHF power, hydrogen dilution, and substrate temperature onthe formation of dust in hydrogen diluted silane plasmas. We used substratetemperatures Ts of 100°C, 150°C, and 200°C, which were reached through res-


0 5 1 0 1 5 2 0 2 5 3 00 . 4 00 . 4 50 . 5 00 . 5 50 . 6 00 . 6 50 . 7 0

H b/H a


γ ’ - r e g i m e α - r e g i m e

Figure 3.4: The intensity ratio Hβ/Hα as an indication of electron temperatureas a function of position between the electrodes in the α-regime (circles) andin the γ’-regime (squares). The position is measured from the upper groundedelectrode downwards.


istive heating of the substrate holder. Although some heating of the substratedue to power dissipation from the plasma is expected, measurements haveshown the additional heating to be less than 3°C for a 20W plasma at typicaldeposition conditions. Because all powers used in these experiments are wellbelow 20W, the plasma heating does not have a significant contribution to thesubstrate temperature.

For all plasmas we used a process pressure of 1.1 mbar, an excitation fre-quency of 60 MHz, a total P ranging from 5 to 20W and hydrogen dilutionR ranging from 20 to 60. The hydrogen flow was kept constant at 100 sccmand the hydrogen dilution was changed by adjusting the silane flow. The dis-tance between the horizontal powered lower electrode and the upper groundedelectrode was 27mm.

If we fix the plasma parameters; pressure, hydrogen dilution, and substratetemperature, we can control the plasma regime by changing the applied power.If we start a plasma in the α-regime and increase the applied power, eventuallydust particles will start to form and the plasma will transit to the γ’-regimeresulting in the described change in the axial optical emission profile.

Using our optical technique of analysing the asymmetry of OES emissiondistribution, we mapped the transition from the α-regime to the γ’-regime asa function of hydrogen dilution, applied power, and substrate temperature.Figure 3.5 shows the results from these investigations. Apart from appliedpower and hydrogen dilution, we clearly see that the transition depends onsubstrate temperature, going into the dusty regime at higher hydrogen dilutionor lower power at lower substrate temperatures. This implies that depositing atlow substrate temperature limits the parameter space for dust-free deposition.Also shown is the transition from amorphous to nanocrystalline growth, whichmainly depends on hydrogen dilution. Together, the two transitions define aparameter window in which we can grow dust-free amorphous silicon, which isvery limited at low substrate temperatures. The parameter window for dust-free amorphous silicon growth at a substrate temperature of 100° is indicatedin grey in the figure. Similar windows can be identified for depositions athigher substrate temperatures. As the deposition rate is directly related tothe power dissipated to the plasma, limitations on applied power will limit themaximum achievable deposition rate.

3.3.3 TEM images of dustIn both regimes we deposited amorphous silicon layers to investigate the pres-ence of dust in the deposited layers. The layers were deposited on a 1µm thickZnO:Al(0.5%) layer on a glass substrate. The ZnO layer was removed by chem-


2 0 3 0 4 0 5 0 6 0 7 00

5

1 0

1 5

2 0

Forw

ard Po

wer (W

)

a - S i n c - S i

Figure 3.5: Map of the transition from the α-regime to the γ’-regime as afunction of hydrogen dilution and applied power at substrate temperaturesof 200°C (squares), 150°C (circles) and 100°C (triangles). The amorphousto nanocrystalline transition also depends on applied power and hydrogendilution. The parameter window for dust free amorphous silicon growth at100°C is shown in grey.

Figure 3.6: TEM image of an a-Si layer deposited in the α-regime (a) and ofan a-Si layer deposited in the γ’-regime (b,c), as identified by analysis of theoptical emission profile. Image (c) shows a part of image (b) in more detail.


ical etching in a 1.5% HCl solution for 1 hour, which leaves flakes of siliconfloating in the solution. This solution was filtered and afterwards washed withethanol. Finally, a transmission electron microscopy (TEM) grid was used toscoop flakes from the ethanol and used as a sample in the TEM microscope.Figure 3.6 shows TEM images of a layer deposited in the α-regime (a) and alayer deposited in the γ’-regime (b and c). In the layer deposited in the γ’-regime we observe particles with sizes ranging from several tens of nanometresup to micrometers, whereas in the layers deposited in the α-regime we observeno particles in the layers. The transition from the α-regime to the γ’-regimewas induced by a small increase in the applied power into the plasma, whereasall the other plasma parameters were kept constant.

3.3.4 OES of pulsed PlasmasA known method to suppress powder (dust) formation is the use of amplitudemodulated plasmas [83, 84, 85]. In contrast to continuous wave (CW) plasmas,the VHF input signal is modulated by a square wave [86]. The behaviourof modulated plasmas is quite different from CW driven plasmas, which ismanifested in a change in deposition rate and material properties and a changein optical emission from the plasma. In this study, we investigate how theoptical emission from the plasma is influenced when a plasma that is in theγ′-regime is influenced when the power input is pulsed. For this purposewe chose the following plasma parameters: Gas flows of 35 sccm SiH4 and175 sccm H2, a time-average applied power 10W, a VHF frequency of 60MHz,an electrode distance d of 27mm, a pressure p of 0.6mbar and a substratetemperature of 130°C. The duty cycle was 50% in all cases. The pressurewas adjusted in such a way that, in a continuous wave plasma, the plasmajust crossed the boundary to the powder forming γ’-regime. For this plasmawe pulsed the power supply at different frequencies: 50Hz, 500Hz, 1 kHz,10 kHz and 100 kHz. We recorded the axial emissions of these plasmas usingthe method described in section 3.3.1. Figure 3.7 shows the axial profilesunder these conditions for the SiH* line emission. In the CW case, it shows atypical emission profile from a γ’-plasma, showing an asymmetric profile anda large contribution from the bulk. We do not observe a dip in the emissionfrom between the sheaths, which is due to the lower pressure compared tothe previous measurements. When the plasma is pulsed at 50Hz, 500Hz, or1 kHz, the emission drops drastically over the whole profile, whereas the shapechanges into a typical profile from a plasma in the α-regime, with articulatedemissions from the plasma sheaths and a low bulk emission. Towards highermodulation frequencies, the overall intensity increases and the bulk intensity


0 5 1 0 1 5 2 0 2 5 3 005

1 01 52 02 5

SiH* In

tensit

y (a.u

.)

V e r t i c a l P o s i t i o n ( m m )

C W P u l s e d 5 0 H z P u l s e d 5 0 0 H z P u l s e d 1 k H z P u l s e d 1 0 k H z P u l s e d 1 0 0 k H z

Figure 3.7: Vertical optical emission profile of the SiH* line of continuous waveand pulsed plasmas. The CW wave plasma was in the γ’-regime. Pulsing theplasma changes the optical emission. Deposition on the window was minimizedby closing the shutter in between measurements.

also increases. Other studies have shown that layers grown in pulsed plasmasshow large amounts of small particles and that their size can be controlledby changing the duty cycle of the modulation [87, 88, 89]. A longer dutycycle will increase the size of the particles. If during the plasma off-period theelectric field within the plasma collapses, the particles can escape the plasmazone if their typical diffusion time is shorter than the plasma off time. Thediffusion time depends on a number of plasma properties such as pressureand temperature, and on the reactor geometry. Judging from figure 3.7, weestimate the typical diffusion time to be between 0.05 and 0.5ms in our reactorunder these specific plasma conditions.

3.4 Mass spectrometry3.4.1 Clusters, the precursors of dust formationIt has been shown in earlier reports that in a silane plasma the energy distri-bution of ions that reach the substrate depends on the substrate temperature

3.4. Mass spectrometry 43

[38]. It is believed that the transfer of energy to the substrate by energeticions plays an important role in obtaining good quality materials [90]. Thishypothesis motivated us to study the substrate temperature dependence ofthe ion energy distribution function (IEDF) of a silane/hydrogen plasma as afunction of substrate temperature. Using a Hiden EQP 300 energy-selectivemass resolved spectrometer we measured the IEDF of several species reachingthe growing surface, through a 30µm sampling orifice fixed at the groundedelectrode, where in a deposition chamber the growing surface would be located.For a fixed gas flow ratio ΦH2/ΦH2 of 5, an applied VHF power of 10W andan excitation frequency of 50MHz we measured the IEDFs of the SinH+

2n+1(n = 1. . . 5) and of H+

2 ions as a function of substrate temperature at differ-ent reaction chamber pressures from 0.05 to 0.25mbar. This parameter spacecorresponds to conditions we use for depositing silicon films for amorphoussilicon solar cells. At these process settings, the plasma is in the dust freeα -regime when the temperature is at the optimum deposition condition ofaround 200 °C, but at low temperature and high pressure settings it is closeto the transition to the dusty γ’-regime. Because larger polysilane ions playan important role in the initial phase of dust formation, the detection of theseions can contribute to the understanding of the temperature dependence ofthe regime change from the α to the γ’-regime [64, 82].

3.4.2 Ion energiesA summary of our results for the ion energies of SiH+

3 (mass = 31 amu) ispresented in figure 3.8 as a function of substrate temperature at differentpressures. In the pressure range studied, the ion energies for all measuredspecies show an increasing trend with increasing substrate temperature. Thistrend is also observed for all the other measured ions in this study, includingH+

2 , which can be ascribed to a longer mean free path at a higher substratetemperature, by complying to the ideal gas law: When the pressure is keptconstant, the gas density decreases when the gas temperature is increased.

3.4.3 Cluster formation and temperatureFigures 3.9 shows the count rate for SiH+

3 at the peak ion energy position asa function of substrate temperature at different chamber pressures. Figure3.10 shows the equivalent for Si4H+

9 ions. The SiH+3 ions’ count rate shows

an increasing trend with increasing temperature, whereas Si4H+9 shows an ini-

tial increase in maximum count rate with increasing temperature region toreach a maximum at low substrate temperature and thereafter, a decrease in


0 5 0 1 0 0 1 5 0 2 0 0 2 5 0

1 5

2 0

2 5

Peak

Ion E

nergy

(eV)

p ( m b a r ) 0 . 0 5 0 . 1 0 0 . 1 5 0 . 2 0 0 . 2 5

S i H 3+

Figure 3.8: The peak ion energy for SiH+3 ions in a hydrogen diluted silane

plasma as a function of substrate temperature at various pressures. The ionenergy rises with increasing substrate temperature and decreases with increas-ing pressure.


0 5 0 1 0 0 1 5 0 2 0 0 2 5 002 0 0 0 04 0 0 0 06 0 0 0 08 0 0 0 0

1 0 0 0 0 0 S i H 3+

Coun

t rate

(C/s)

p ( m b a r ) 0 . 0 5 0 . 1 0 0 . 1 5 0 . 2 0 0 . 2 5

Figure 3.9: The count rate of SiH+3 ions in a hydrogen diluted silane plasma

at the peak ion energy as a function of substrate temperature at differentpressures.

count rate at increasing temperature. For Si5H+11 ions (not shown) the count

rates decrease monotonically with increasing temperature for all pressures. Wespeculate that the increase of the ion energies with substrate temperature cor-responds to the larger mean free path at higher gas temperatures. Becauseof this, the positive ions have a higher probability of travelling through theplasma sheath near the grounded electrode without colliding; their energieswill be higher and so will be their count rate. Decreasing the pressure willincrease the mean free path of the clusters and will therefore have a similareffect. Due to the increase in mean free path, the ion energy and the countrate for low-mass ions increases, as observed for silyl ions. It is known thata higher gas temperature reduces the polymerization rate of silyl ions intolarger polysilanes [64]. The increase in the number of positive ions throughthe plasma sheath is counteracted by the decrease in polymerization rate, asobserved at higher temperatures. Figure 3.11 again shows count rates as afunction of pressure and temperature, but now for Si3H+

7 ions. We observetwo opposing trends: At low pressures (0.05 and 0.10mbar) the increase insubstrate temperature suppresses the formation of clusters, whereas at higherpressures (0.15, 0.20, and 0.25 mbar), due to the larger number of collisions


0 5 0 1 0 0 1 5 0 2 0 0 2 5 00

5 0 0

1 0 0 0

1 5 0 0S i 4 H 9

+

Coun

t rate

(C/s)

p ( m b a r ) 0 . 0 5 0 . 1 0 0 . 1 5 0 . 2 0 0 . 2 5

Figure 3.10: The count rate of Si4H+9 ions in a hydrogen diluted silane plasma

at the peak energy as a function of substrate temperature at different pressures.

at elevated pressures, the trend towards higher substrate temperature is stilldominated by the increase in mean free path. It should be noted that we areonly able to measure positive ions. Because positively charged particles areeasily ejected from the plasma due to the plasma potential profile, the clustersunder investigation formed through two possible routes can be detected: ifthe polymerization occurs fast, i.e. within the typical time needed for a pos-itive particle to be ejected from the plasma, the cluster can spawn throughpolymerization of positive particles by insertion of SiH2 into SinH

+m ions,

although this reaction is believed to stop above n = 6 or 7 [67]. The other pos-itive cluster generation scheme is through the positive charging of negativelycharged clusters, which can happen through the collision with a positivelycharged ion (anion-cation neutralization [67]), making it a neutral particle,followed by a discharge reaction (SinH2n+2 + e− → SinH

+2n+1 + H + 2e−),

ionizing the particle to a positively charged ion. Because the polymerizationof negative polysilanes occurs faster than the polymerization of positive chains[91], the latter is the most probable route towards positively charged highersilanes. Our results, showing a decrease in count rate for higher mass ions withincreasing temperature, confirms a negative influence of the gas temperature


0 5 0 1 0 0 1 5 0 2 0 0 2 5 01 0 0 0

2 0 0 0

3 0 0 0

4 0 0 0 S i 3 H 7+ p ( m b a r ) 0 . 0 5

0 . 1 0 0 . 1 5 0 . 2 0 0 . 2 5

Coun

t rate

(C/s)

Figure 3.11: The count rate of Si3H+7 ions in a hydrogen diluted silane plasma

at the peak energy as a function of substrate temperature at different pressures.The temperature dependence of the cluster growth changes with pressure.

on cluster formation. This confirms also the hypothesis that a higher substratetemperature will prevent the plasma from going into the dusty γ’-regime bysuppressing the polymerization reactions of silyl into larger polysilanes.

3.4.4 ConclusionsWe have presented a non-invasive in-situ technique to determine whether adeposition plasma is in the dust free α-regime or in the dusty γ’-regime byrecording a spatially resolved optical emission profile perpendicular to theelectrodes. In the γ‘-regime we observe an increase in electron temperature inthe bulk, which indirectly confirms the presence of dust particles. We mappedthe transition from the α- to the γ’-regime as a function of hydrogen dilu-tion, applied power and, substrate temperature. This method can generallybe used to determine the processing window in which dust free silicon filmscan be deposited. Because optical emission can also be used to predict thematerial phase, this study shows that a single technique can be used as anin − situ plasma diagnosis tool for monitoring of the amorphous to nano-crystalline transition as well as the transition from the dust-free to the dusty


regime without a supplement technique. Therefore it is a plasma monitoringtool to control the complete thin film silicon processing of ’micromorph’ cellmanufacturing.

Dust formation can be suppressed by using a pulsed power input. Changingthe modulation frequency and thus changing the plasma-off time determineswhether dust formation is suppressed.

We found a dependence of the initial growth of positively charged clusterson substrate temperature, which explains the temperature dependence of theα to γ’ transition.

Chapter 4

Low temperature siliconlayers

4.1 The role of substrate temperature in PECVDThe electronic quality of thin amorphous silicon (a-Si:H) and nanocrystallinesilicon (nc-Si:H) films is directly influenced by their deposition temperature.For both materials a minimum in defect density (dangling bonds) is observedfor layers grown at a substrate temperature of around 200°C to 250°C [36].Dangling bonds are unoccupied silicon bonds and these can act as recombin-ation sites for charge carriers, but can be passivated by the attachment of ahydrogen atom. Deposition at lower than standard substrate temperatureswill induce a higher defect density as well as a higher porosity and thus lowerrefractive index [92, 93], and therefore, lead to lower solar cell efficiencies.

An a-Si:H deposition at standard temperature (~200°C to 250°) goes throughthermal equilibrium [94], which allows for minimisation of defects and optimumhydrogen content for dangling bond passivation. In this thermal equilibrium,there is a balance between strained intersilicon bonds (Si—H—Si) and Si—Hbonds: Si—H + Si——Si⇐⇒ Si—+ Si—H—Si [95]. Below standard tem-perature, which is associated with the glass transition temperature of a-Si:H,the structure is ’frozen’, and the hydrogen cannot move around to passivatedefect sites. In this non-equilibrium regime, reactions in the growth zone ofthe growing film are important. Passivation occurs when growth precursors(mainly SiH3 [36]) move across the growing surface. Therefore the defect dens-ity is related to the diffusion length of the precursors, which is related to the

50 Chapter 4. Low temperature silicon layers

substrate temperature [96].The surface mobility and therefore the material quality can be improved

by adding hydrogen to the feedstock gasses, so-called hydrogen dilution. Bydoing this, even at a low substrate temperatures of 100°C or 75°C workingdevices can be fabricated [97, 98]. Although the quality of low temperaturesilicon layers can be improved by increasing the hydrogen dilution, doing sowill generally decrease the deposition rate .

It is generally accepted that energy transported to the growing surfaceby ion bombardment in the plasma enhanced chemical vapour deposition(PECVD) process can contribute to a higher surface mobility of these mo-lecules. The kinetic energy flux carried by the energetic ions can be variedby changing the plasma properties, such as the applied power density, the gaspressure or the hydrogen dilution. It has also been shown that the ion en-ergy flux towards the substrate decreases when the substrate temperature isdecreased, which in turn can be increased by adding more hydrogen. It hasbeen shown that, going from a deposition at 200°C to a deposition at 39°C,the energy flux towards the growing surface can be restored to the originallevel when the hydrogen dilution is increased [38].

Changing the substrate temperature also changes the optical properties ofthe silicon. Figure 4.1 depicts the absorption coefficient and the refractive in-dex as a function of wavelength of typical a-Si:H layers deposited at substratetemperatures of 180°C and 130°C as measured by reflection-transmission meas-urements (R-T) and fitted to the OJL-model [55], showing a lower refractiveindex over the whole spectrum and a lower absorption coefficient for the layerdeposited at low substrate temperature. These optical properties clearly in-dicate a higher bandgap for a-Si:H deposited at lower temperatures, which canbe attributed to a higher hydrogen content in the layers [92].

This chapter will cover our search for device quality intrinsic a-Si:H andnc-Si:H layers, as well as p-type and n-type doped layers to be used for de-positions directly onto polycarbonate (PC), which limits the substrate tem-perature to around 130-140°C, due to the glass transition temperature of PC[34].

4.2 Controlling the substrate temperatureWhen depositing on plastic substrates, it is very important to accurately mon-itor and control the substrate temperature. A temperature higher than theglass transition temperature (Tg) of the substrate will result in deformationor even melting of the substrate. Furthermore, in the range from 100°C to

4.2. Controlling the substrate temperature 51

4 0 0 5 0 0 6 0 0 7 0 0 8 0 0234

1 0 1

1 0 3

1 0 5 Ab

s. Co

eff. (c

m-1 )


Refr.

Index

Figure 4.1: Refractive index and absorption coefficient of a-Si:H deposited at130°C and at 180°C and of sputter deposited ZnO:Al.

160°C, the silicon material quality is very sensitive to small variations in tem-perature [36] and therefore it is important to establish a controlled and stabletemperature.

4.2.1 Substrate stretch holderWhen a flexible substrate is fixed to a rigid substrate holder and the thermalexpansion coefficient of the substrate material is higher than that of the sub-strate holder, the substrate will curve when it expands due to thermal ex-pansion during deposition. This is the case for everyday plastics, which havethermal expansion coefficients much higher than of titanium, which is the ma-terial our substrate holders are made of. This curving will cause a gap betweenthe substrate and the substrate holder, which will result in a suboptimal andan inhomogeneous heat transfer from the substrate holder to the substrate,resulting in a inhomogeneous and lower quality layer. Furthermore, becauseintroducing a gap between the substrate and substrate holder locally changesthe electrical properties of the plasma, a change in deposition rate occurs [99].Therefore curving of the substrate should be avoided. To do this, we useda specially designed substrate holder that can stretch the substrate when it


Figure 4.2: The substrate stretch holder, which is used to maintain a goodcontact between the substrate holder and the substrate and to avoid curvingof the substrate. The maximum size of the substrate is 10 ×10 cm2.

expands and thereby keeping it flat. Figure 4.2 shows a photograph of thisstretch holder, with in the inset a detailed image of one of the springs pullingon the substrate. This configuration resembles the situation for a roll-to-rollprocess, in which also the substrate is kept under tension during the deposition[100]. Due to the design, there is always a small gap between the substrate andthe holder. The part of the stretch holder which holds the substrate is madeof stainless steel, whereas the normal substrate holders are completely madeof titanium. The problem of curving substrates is only relevant in the case ofbatch-type processing, whereas in roll-to-roll processing we do not expect thisproblem.

4.2.2 Gas pressureTo calibrate the actual temperature of the substrate to the accurately con-trolled heater temperature, we mounted both a glass substrate and a PC sub-strate to a regular and to the stretch substrate holder. To both substrateswe attached K-type (chromel-alumel) thermocouples in the centre and at 1 cmfrom the edge of the substrate. To simulate deposition conditions we intro-duced 0.16mbar or 5mbar argon gas into the reactor and ignited a plasmarunning at 5 or 20Watts to mimic a-Si:H and nc-Si:H deposition conditions,


1 0 0 1 2 5 1 5 0 1 7 5 2 0 05 07 5

1 0 01 2 51 5 01 7 52 0 0

G l a s s , v a c u u m G l a s s , 0 . 1 6 m b a r A r G l a s s , 5 m b a r A r P C , 0 . 1 6 m b a r A r P C , 5 m b a r A r 5 / 7 r u l e

Figure 4.3: The substrate temperature measured on glass substrates in a reg-ular substrate holder and on PC mounted on the stretch substrate holder,as a function of set heater temperature at two different argon gas pressures,corresponding to a-Si:H and nc-Si:H deposition pressure conditions.

respectively. For these configurations we measured the relation of the substratetemperature to the heater temperature in the IRIS deposition chamber. As arule of thumb, we normally estimate the substrate temperature to be 5/7 ofthe set heater temperature. The measurements were performed in a separatedeposition setup called IRIS, which uses an excitation frequency of 50MHz,whereas intrinsic layers in the ASTER system are deposited at 60MHz and thepowered electrode is not a showerhead electrode, in contrast to the ASTERsetup. Otherwise the IRIS reactor chamber is a near-exact copy of one of theASTER deposition chambers.

Because there can only be a limited number of contact points between theheater and the substrate holder and between the substrate holder and the sub-strate, most of the heat transfer (in high vacuum conditions) is of a radiativenature and therefore the size of the gap between the heater and the substrateholder and between the substrate holder and the substrate is important forthe heating of the substrate. Introducing gas into the reactor will induce


convective heat transfer, causing a higher substrate temperature at a fixedheater temperature. At the end the substrate temperature will depend on theproperties of the substrate material, i.e. thermal conductivity and emissivity.Figure 4.3 shows the temperature calibration results for a glass substrate ina regular substrate holder and a PC substrate in the stretch holder. Meas-urements were performed under vacuum condition and in a 0.16mbar and a5mbar argon environment with no plasma running. The same pressures areused for a-Si:H or nc-Si:H depositions, respectively. Going from vacuum con-ditions to a low pressure condition such as 0.16mbar of argon (although amixture of silane and hydrogen is used in silicon film depositions), the sub-strate temperature increases by a few degrees. A small difference in substratetemperature is observed between the glass substrate in a regular substrateholder and a PC substrate in the stretch holder, which could be attributedto a small gap between the holder and substrate due to the holder design, ordue to the fact that the stretch part of the holder is made from stainless steel,whereas the regular holder is entirely made of titanium. Titanium has a muchhigher thermal conductivity coefficient (21.9 Wm−1K−1 [101]) than stainlesssteel (12-14 Wm−1K−1 [102]).

4.2.3 Plasma heatingRunning a plasma inside the reactor will raise the substrate temperature.Coupling power into the plasma will induce heating of the substrate [103]. En-ergy from the plasma is transferred to the substrate through ion bombardment.Figure 4.4 shows temperature calibration measurements of a glass substrateand a PC substrate, both mounted in the stretch substrate holder, under twodifferent plasma conditions. Before the measurements we waited for the tem-perature to stabilize at a given pressure. The top graph shows the heatingof the substrate as a function of time under our most-used low temperaturea-Si:H deposition conditions, whereas the bottom graph shows the same forlow temperature deposition conditions for device quality nc-Si:H. The plasmasettings are listed in table 4.1. For both plasma regimes, the substrate heatsup during the deposition and the glass substrate temperature is always 2-3°Chigher than the PC substrate temperature. In both regimes, during the firstfew minutes there is a rapid increase in substrate temperature, which graduallychanges into a linear temperature increase of around 1°C in 3.5 minutes. Thelinear part of the heating is probably due to heating of the complete reactorchamber. During a typical 300 nm a-Si:H i-layer deposition ( 25 minutes) theheating due to the plasma is 7-8°C. During a 1000 nm nc-Si:H i-layer deposition( 30 minutes) the plasma heating is estimated to be around 11°C.


0 2 4 6 8 1 01 3 01 3 51 4 01 4 5 0 5 1 0 1 5 2 0 2 51 1 51 2 01 2 51 3 0

P l a s m a - o n t i m e ( m i n u t e s )

G l a s s P C

1 7 . 5 W a t t3 . 0 m b a r

G l a s s P C

5 . 0 W a t t0 . 1 6 m b a r

Figure 4.4: The substrate temperature measured on a glass substrate and on aPC substrate, both mounted on the stretch substrate holder, as a function ofplasma-on time. The plasma conditions are that of a standard low temperaturea-Si:H growing plasma and a nc-Si:H growing plasma.

Plasma p P ΦSiH4 ΦH2 d Thmbar W sccm sccm mm °C

a-Si:H 0.16 5.0 35 175 27 180nc-Si:H 3.0 17.5 5 100 10 170

p: pressure; P: applied plasma power; Φ: gas flow; d: interelectrode distance; Th: heater temper-ature

Table 4.1: Plasma properties as used for temperature calibrations in IRIS. Thecalibration results are shown in figure 4.4.


1 0 0 1 1 0 1 2 0 1 3 0 1 4 0 1 5 0

1 0 1 7

2 x 1 0 1 7

3 x 1 0 1 74 x 1 0 1 75 x 1 0 1 76 x 1 0 1 7

D e f e c t D e n s i t y

Defec

t Den

sity (

cm-3 )

5 0

5 5

6 0

6 5

7 0

U r b a c h E n e r g y Urba

ch En

ergy (

meV)

Figure 4.5: Midgap defect density and Urbach energy as a function of substratetemperature for a-Si:H layers, derived from CPM measurements.

4.3 Low temperature intrinsic layers

4.3.1 a-Si:H intrinsic layers

Temperature series

A series of a-Si:H layers was deposited on glass substrates with substrate tem-peratures ranging from 100°C to 145°C to investigate the influence of thesubstrate temperature on the layer quality within our temperature range ofinterest. The deposition parameters were: ΦSiH4 : 35 sccm, ΦH2 : 175 sccm,p: 0.16mbar, P: 5W, d: 27mm. Layer thicknesses were around 600 nm. Thetemperature was varied by changing the controlled heater temperature andapplying the 5/7-rule. No large differences are found in band gap or light anddark conductivities. However, we do find a clear trend in midgap defect dens-ity and Urbach energy (Eu), both derived from constant-photocurrent methodmeasurements (CPM). The results from this study are shown in figure 4.5. Go-ing from low to high temperature the defect density drops from over 5×1017

to 8×1016 cm−3, whereas the Urbach tail energy decreases from 69 to 58 meV.

4.3. Low temperature intrinsic layers 57

1 0 0 1 2 5 1 5 0 1 7 5 2 0 0

2 x 1 0 1 7

4 x 1 0 1 7

6 x 1 0 1 78 x 1 0 1 7

1 0 1 8

D e f e c t D e n s i t y

De

fect D

ensit

y (cm

-3 )

H y d r o g e n F l o w ( s c c m )

5 5

6 0

6 5

7 0

7 5

U r b a c h E n e r g y Urba

ch En

ergy (

meV)

1 0 0 1 2 5 1 5 0 1 7 5 2 0 0 2 2 50 . 1 20 . 1 40 . 1 60 . 1 80 . 2 00 . 2 2

R*

H y d r o g e n F l o w ( s c c m )

Figure 4.6: Defect density and Urbach energy as a function of hydrogen di-lution, derived from CPM measurements for a-Si:H layers deposited at a sub-strate temperature of 130°C. The inset shows the microstructure factor, asobtained from FTIR measurements.

Hydrogen dilution series

To achieve optimal layer quality for device performance, a hydrogen dilutionseries was performed, using a silane flow of 35 sccm, and a varying hydrogengas flows between 105 and 200 sccm. The other deposition parameters were:p: 0.16mbar, P: 5W, d: 27mm, Ts: 130°C. The results for the midgap defectdensity and Urbach energy, obtained from CPM measurements and the micro-structure parameter R*, obtained from fourier transfer infrared (FTIR) dataare presented in figure 4.6. It shows an increasing layer quality trend towardhigher dilutions, expressed in a lower defect density, which changes almost anorder of magnitude while increasing the hydrogen to silane ratio from 3 toaround 6. Also the Urbach energy changes from more than 70meV at the lowdilution end to 56meV for the layers grown at the highest dilution. Becausealso the deposition rate changes only a few percent, layers grown at high dilu-tion would be most appropriate for device production. However, the highestdilution ratio layers suffer from high compressive stress, resulting in peelingof the layers from the substrate. It is known that a rise in compressive stress


occurs just before the transition from a-Si:H to nc-Si:H, which is attributed toa higher hydrogen content in the layers, which partially is molecular hydrogentrapped in microvoids that contributes to macroscopic stress [104].

4.3.2 nc-Si:H intrinsic layersIn the nc-Si:H growth regime, we performed a hydrogen dilution series anda plasma-power series to find an optimal recipe for depositing intrinsic nc-Si:H for solar cells, both in single-junction cells as well as for the bottomcells of a-Si:H/nc-Si:H tandem cells. Because our aim is to deposit theselayers on polycarbonate substrates, the substrate temperature of these layersis kept at 130°. The key physical parameter for these layers is the crystallinefraction, which can be determined by Raman-spectroscopy. The procedurefor this is described in section 2.2.3. Silicon films grown near the transitionregime, consisting of a Raman crystalline ratio from 0.3 to 0.5 are referred toas mixed-phase or transition-type silicon and are the preferred materials fornc-Si:H cells and a-Si:H/nc-Si:H solar cells [105, 106]. Because these layers aregrown under plasma conditions for which the nucleation rate of crystallites islow, small changes in the plasma or in the substrate material composition ormorphology can change the crystalline fraction, making it difficult for theselayers to reproduce. The use of a crystalline seed layer will solve this issue,because nucleation has already taken place. Furthermore, if these layers aregrown in a p-i-n solar cell structure, they will be deposited on a nc-Si:H p-typelayer, which acts as a seed layer. Using such a p-layer as a seed layer wouldrender it useless for conductivity measurements of the layers. To mimic thelayer growth as if it was part of a solar cell structure, these layers were grownon a thin (∼10 nm) nc-Si:H seed layer, which was deposited from a nc-Si:Hp-layer recipe, but without adding the trimethylborane to the gas mixture,which is used to dope the material p-type. This is done so that we are stillable to do conductivity measurements. If we were to use a doped nc-Si:H seedlayer, the conductivity would be dominated by the conductivity of the p-layer.

Dilution series

Figure 4.7 shows the crystalline fraction and deposition rate of layers grownat different hydrogen dilutions (while keeping the H2 flow fixed) around thetransition from a-Si:H to nc-Si:H growth. As expected, we observe a transitionfrom a-Si:H to nc-Si:H when the silane concentration in the gas feedstock mix-ture is increased. To evaluate whether the crystalline fraction is homogeneousthroughout the layer in the growth direction, the Raman spectra were meas-

4.3. Low temperature intrinsic layers 59

4 5 6 70 . 0

0 . 2

0 . 4

0 . 6

0 . 8

1 . 0 C r y s t a l l i n e f r a c t i o n L a y e r S i d e G l a s s S i d e

Cr

ystal

line R

atio

S i l a n e F l o w ( s c c m )

0 . 5

0 . 6

0 . 7

0 . 8

0 . 9

Dep

ositio

n Rate

(nm/

s)

D e p o s i t i o n R a t e

Figure 4.7: Crystalline fraction and deposition rate of Si layers deposited at130°C on a nc-Si:H seed layer as a function of silane content of the plasmafeedstock mixture, measured by R-T and Raman-spectroscopy.

ured both from the top of the layer and through the Corning glass substrate.This study reveals a comparable crystal fraction on both sides, indicating ho-mogeneous growth in the growth direction. Although the measured crystalfraction from the substrate side is influenced by the presence of the nc-Si:Hseed layer, the nc-Si:H seed layer only has a minor influence on the measuredcrystalline fraction, because the penetration depth for Raman measurements isaround 100 nm for nc-Si:H, when using laser light with a wavelength of 514 nm[107], whereas the seed layer is only 20 nm. The deposition rate decreases withincreasing hydrogen dilution. The hydrogen flow was 100 sccm, while the otherplasma parameters were p: 3.0mbar, P: 20W, d: 10mm, Ts: 130°C. The layerthickness was around 800 nm.

The desired Raman crystalline ratio is found at a silane flow of around5 sccm. This layer is used for further optimalisation of the intrinsic nc-Si:Hlayers for use in a-Si:H/nc-Si:H tandem structures. This is done by varyingthe applied power into the plasma.


1 0 1 5 2 0 2 5 3 00 . 0

0 . 2

0 . 4

0 . 6

0 . 8

1 . 0

C r y s t a l l i n e f r a c t i o n L a y e r S i d e G l a s s S i d e

Crys

tallin

e Rati

o

P o w e r ( W )

0 . 3 50 . 4 00 . 4 50 . 5 00 . 5 50 . 6 00 . 6 5

D e p o s i t i o n r a t e

Dep

ositio

n Rate

(nm/

s)

Figure 4.8: Crystalline fraction and deposition rate, measured by R-T andRaman-spectroscopy, of Si layers deposited at 130°C on a nc-Si:H seed layeras a function of applied power into the plasma.

Power series

The nc-Si:H layers developed from the hydrogen dilution series was used asa basis for fine-tuning of the crystal fraction by changing the applied powerinto the plasma. The power input was varied between 10 and 30W in 5Wincrements. The gas flows were: silane: 5 sccm, hydrogen: 100 sccm. Theother plasma parameters were p: 3.0mbar, d: 10mm, Ts: 130°C. The layerthickness was around 800 nm. The results for the Raman crystalline ratio,measured both on the layer top surface and through the glass substrate aswell as the deposition rate are presented in figure 4.8. In this regime a trans-ition from completely amorphous to a Raman-crystalline fraction of 80% takesplace. Between 15W and 20W, the crystal fraction changes from ∼25% to∼50%, which is in the desired mixed-phase regime. For powers up to 20W, thedeposition rate increases with increased power input but saturates at higherapplied powers, indicating silane depletion conditions. The layers deposited inthis power series were chosen as the i-layer material for nc-Si:H single junctionand a-Si:H/nc-Si:H tandem solar cells, as described in chapter 6.

4.4. Low temperature doped layers 61

Description Ts ΦSiH4 ΦH2 ΦCH4 ΦT MB ΦP H3 P p°C sccm sccm sccm sccm sccm W mbar

a-Si:H p 130 35 124 - 53 - 5 0.16a-SiC:H p 200 30 - 40 18 - 5 0.15a-Si:H n 130 30 23 - - 7.5 5 0.5a-Si:H n 200 30 23 - 7.5 5 0.5nc-Si:H p 130 1 240 - 0.25 - 15 1.1nc-Si:H n 130 1.2 180 - - 0.27 15 1.1

Table 4.2: Gas flows, applied power and pressure used for plasma depositionof doped layers at substrate temperatures of 130°C and 200°C. The measuredproperties of these layers are summarized in table 4.3. TMB and PH3 arediluted in H2 gas at a concentration of 2 at. %.

Nanocrystalline layers on PC

Because the nucleation of crystallites can be sensitive to the substrate material,investigations have been performed on the crystal growth on PC substrates.The results are presented in chapter 6.

4.4 Low temperature doped layersDoped layers were deposited at a substrate temperature of 130°C by adaptingthe substrate temperature for known optimised recipes at higher (standard)temperatures or from optimized layers developed previously for solar cells at100°C for deposition in the ASTER deposition system [100], and character-ized. For this purpose p-type a-Si:H, n-type a-Si:H, p-type nc-Si:H and n-typenc-Si:H were deposited and characterized. Table 4.2 shows the gas flows, ap-plied power and reactor gas pressure of these depositions. All doped layerswere deposited at a plasma frequency of 50MHz, an interelectrode distance of27mm and a substrate temperature of 130°C. Table 4.3 shows the measuredoptical and electrical properties of the deposited layers, which were used forlow temperature silicon thin film depositions as used for single junction a-Si:Hand nc-Si:H cells and a-Si:H/nc-Si:H tandem solar cells.

The a-Si:H doped layers at 130°C show electrical properties very similar toa-Si:H doped layers deposited at 200°C in our lab. Optically, the band gapsof a-SiC:H at high temperature are higher due to the incorporation of carbonin the network. For the nc-Si:H doped layers we have no information availableon the standard temperature (200°C) counterparts to compare with.


Description Ts rd Eg(Tauc) Eg(cubic) E04 Ea σd Rc

°C nm/s eV eV eV eV (Ωcm)−1 %

a-Si:H p 130 0.244 2.04 1.69 2.03 0.41 2.3×10−6 -a-SiC:H p 200 0.188 2.10 1.75 2.09 0.45 7.7×10−7 -a-Si:H n 130 0.052 1.82 1.46 1.85 0.23 3.0 ×10−3 -a-Si:H n 200 0.08 1.93 1.57 n/a 0.20 9.8×10−3 -nc-Si:H p 130 0.052 2.09 1.75 1.95 0.05 7.8×10−1 57nc-Si:H n 130 0.036 2.35 2.13 2.45 0.06 8.8×10−2 71

Ts: substrate temperature; rd: deposition rate; Eg : band gap; Ea: activation energy; σd: darkconductivity; Rc: Raman crystalline ratio

Table 4.3: Optical and electrical properties of doped silicon thin films depositedat a substrate temperatures of 130°C and 200°C. The deposition parametersof these layers are given in table 4.2.

The insertion of a thin (5 to 10 nm) nc-Si:H p-layer between the ZnO:AlTCO and the a-Si:H p-layer improves the contact between the (n-type) ZnO:AlTCO and the p-type silicon layers [108], resulting in a higher Voc and FF. Be-cause this layer is extremely thin, nucleation of crystallites needs to occurwithin the first few nanometres of layer growth. To achieve this, we aim to de-velop thin, but highly crystalline nc-Si:H p-type doped layers. For this purposewe performed a hydrogen dilution series, using hydrogen flows ranging from150 to 270 sccm. The other plasma parameters were identical to the paramet-ers listed in table 4.2 for the p-type nc-Si:H deposition. The Raman crystallineratio and deposition ratio for these layers are presented in figure 4.9. We findan increasing Raman crystalline ratio with increasing hydrogen dilution. Alllayers in this series show a similar activation energy of around 0.05 eV. Thedark conductivity at 300K of the layers ranges from 3.1 ×10−1 (Ω cm)−1 at ahydrogen flow of 150 sccm to 7.8 ×10−1 (Ω cm)−1 at 240 sccm H2. The layerdeposited at a hydrogen flow of 270 sccm showed very inhomogeneous growthand therefore we could not measure the dark conductivity. Based on theseproperties we decided to use a hydrogen flow of 240 sccm H2 for the nc-Si:Hp-layer for solar cell growth.

4.5 ConclusionsTo be able to deposit silicon thin films on plastic substrates without deform-ing them, the substrate temperature was calibrated for different substratesand two different substrate holders. The deposition chamber gas pressure has

4.5. Conclusions 63

1 4 0 1 6 0 1 8 0 2 0 0 2 2 0 2 4 0 2 6 0 2 8 00 . 3 00 . 3 50 . 4 00 . 4 50 . 5 00 . 5 50 . 6 00 . 6 50 . 7 00 . 7 5

Crys

tallin

e Frac

tion

H y d r o g e n F l o w ( s c c m )0 . 0 3 00 . 0 3 50 . 0 4 00 . 0 4 50 . 0 5 00 . 0 5 50 . 0 6 00 . 0 6 50 . 0 7 00 . 0 7 5

Depo

sition

Rate

(nm/

s)

Figure 4.9: Raman crystalline ratio, measured by Raman spectroscopy anddeposition rate, measured by R-T spectroscopy, of p-type nc-Si:H layers de-posited at 130°C on ZnO:Al coated glass substrates, as a function of hydrogenflow.


a large influence on the substrate temperature. There is also a differencebetween the substrate temperature of glass substrates in a regular substrateholder compared to PC substrates in the substrate stretch holder. Running aplasma will heat up the substrate 7°C to 8°C in a typical 300 nm a-Si:H i-layerdeposition. During a typical 1µm nc-Si:H i-layer deposition, the additionalheating of the substrate is estimated to be 11°C.

We optimized both intrinsic layers and doped layers for use in solar cellsdeposited at a substrate temperature of 130°. In the amorphous regime, i-layers show a monotonous increase in layer quality with increasing hydrogendilution, but at high hydrogen dilutions stress becomes a problem, and shouldtherefore be avoided. In the nc-Si:H regime we can control the crystal fractionof layers both by the hydrogen dilution as well as the applied power into theplasma. In this way we can control the crystallinity of the layers to a desiredvalue.

For doped layers we adapted known recipes at higher and lower substratetemperatures to our desired substrate temperature and found layer character-istics suitable for solar cell depositions.

Chapter 5

Light trapping inamorphous silicon cells onpolycarbonate substrates

5.1 Light trapping techniques

In order to develop highly efficient thin film silicon solar cells in general, andon plastics in particular, light management is a key feature to optimize theabsorption of light over the complete solar spectrum to obtain a high currentdensity. Because thick silicon absorber layers are detrimental for the elec-tronic properties of the solar cells due to collapse of the electric field in theintrinsic layer bulk that leads to increased recombination, light managementtechniques are needed that boost the absorption such that an adequate lightinduced current can be generated while using thin absorber layers. For thinfilm silicon solar cells, several methods have been investigated to enhance thelight absorption in the active layers of solar cells. The use of textured frontTCOs to scatter the incoming light in superstrate type of cells and a roughZnO-silver interface at the back reflector (BR) for substrate type of cells hasbeen extensively studied [109, 110]. Other techniques use 2D or 3D gratingson back reflectors in substrate type of cells [48, 111] or gratings implemen-ted in the substrate surface to induce diffraction of the light into the activeareas of the cells in superstrate configured cells [112]. Also research is doneon nanopillar [49] structured 3D back reflectors in thin film silicon cells in a

66 Chapter 5. Light trapping in amorphous silicon cells on polycarbonatesubstrates

p-i-n configuration. Since cost is a very important factor in the success ofa certain type of cell, the light management techniques do not only need tobe successful in light absorption enhancement, but should also be suitable forlarge scale processing. Furthermore, using thinner absorber layers will poten-tially have a number of positive consequences: lower material usage, lowerprocessing times, and therefore higher throughput, improved stability againstlight induced degradation and a higher open-circuit voltage [25, 113, 28].

In this research, we present results obtained for three different light trap-ping techniques, which were used for thin film amorphous silicon solar cellsdeposited at a substrate temperature below 130°C: (1) Scattering, by using atexture-etched TCO front contact, obtained by etching the ZnO:Al, coated onglass in an HCl solution in water; (2) Using regular textures comparable in sizeto the effective wavelength of visible light; (3) We will introduce the conceptof geometric light trapping, which is based on refraction and reflection of lighton structures larger than the wavelength of visible light.

To test our light trapping techniques, we have developed amorphous siliconintrinsic layers at deposition temperatures of 130°C with acceptable electronicquality for use in solar cells. We also developed n-type and p-type dopedamorphous silicon layers at these low deposition temperatures. Using ZnO:Alas a front transparent conductive oxide and as a BR, we fabricated a-Si solarcells on our different structured polycarbonate (PC) substrates. On thesecells we measured current density-voltage (J-V) characteristics and spectralresponse (SR). As a reference we deposited a-Si cells under the same conditionson flat glass/ZnO and on Asahi U-type TCO glass.

5.1.1 ScatteringA traditional way to enhance light absorption in a superstrate cell configura-tion, is to use a randomly textured TCO, like texture-etched aluminium dopedzinc oxide (ZnO:Al) or natively textured materials such as fluorinated tin oxide(SnO2:F) as used for the commercially available Asahi TCO glass. A roughinterface between the TCO and the silicon layers causes the light to scatterinto the optically active silicon layers, causing a longer light path through thesilicon and thereby enhancing the absorption without increasing the thicknessof the silicon layers. If the scattering angle is large enough, the light that isreflected from the back of the cell will be trapped inside the cell through totalinternal reflection [114, 45]. Because also electrically inactive layers like theTCO or p-layer have a relatively high absorption in the blue part of the solarspectrum, a major drawback of using textured interfaces is that the scatteringis most pronounced in the short wavelength region, which also causes substan-

5.1. Light trapping techniques 67

1 μm

Figure 5.1: AFM image of a 1µm ZnO:Al(0.5%) layer, texture etched in a1.5% HCl solution for 10 seconds, used as a front TCO in p-i-n solar cells.

tial absorption in these layers. As the absorption coefficient of the intrinsiclayer material in the red part of the spectrum is low, a large light path increaseis needed to capture a substantial part of the red light.

5.1.2 Nanopyramid periodic structures

Several groups have investigated the use of 1D and 2D periodic structuresas light trapping schemes. These gratings can either be integrated into thefront contact in superstrate (p-i-n) type of cells or used in the back contactin substrate (n-i-p) type of cells. Although the far-field effects of these peri-odic structures are well understood, understanding the near-field interactionof the periodic structures and the incoming light is crucial for optimization ofthe structures. Different theoretical investigations have been done for severaldifferent types of nanostructures. The main parameter in these investigationsconcerns the period (or pitch) of the nanostructures. Numerical simulationsby Haase et al. [115] and Dewan et al. [116] for a nc-Si cell in a p-i-n con-figuration on pyramid-structures predict an optimal pitch of 850 and 700 nm,respectively. Campa et al. [117] and Gomard et al. [118] theoretically invest-


500 nm 500 nm

Figure 5.2: AFM images of the convex nanopyramid structured PC (left) andconcave nanopyramid structured PC (right).

igated the use of rectangular gratings and photonic crystals for a-Si cells in ann-i-p configuration and found optimum pitches of 300 and 400 nm respectively.Ferry et al. [48, 119] and Zhu et al. [120] demonstrate absorption enhance-ment effects of nanocones and nanodomes in n-i-p a-Si solar cells with verythin absorber layers and find optimum periods of 500 and 450 nm. Eisele etal. show reduced reflectance from 1D gratings in a-Si cell-like layer stacks forgratings with 389 and 798 nm gratings in a p-i-n configuration [121].

In this chapter, we present a-Si cells deposited on nanopyramid structuredsurfaces in a p-i-n configuration. As a substrate we used PC, embossed with400 nm pyramids on a square base. We experimented both with inverted pyr-amids (facing inwards, called type I from now) and normal pyramids (facingoutwards, type II). AFM images of the structures used are displayed in figure5.2, showing convex pyramids on the left and concave pyramids on the right.The pyramid base for both structures is 400 nm, which is expected to givemaximum current enhancement for a-Si thin film solar cells, as studied in theliterature.

5.1.3 Geometric light trapping: micropyramid periodicstructures

We present a light trapping scheme based on reflection and refraction, whichdoes not rely on scattering on nanostructured interfaces. We fabricated suchstructures on PC substrates, by embossing them with 8 µm base pyramids


Figure 5.3: SEM images of convex pyramids (left) and concave pyramids(right) embossed on polycarbonate. These substrates are used for geomet-ric light trapping.

with positive pyramid angles, using a hot embossing technique [122]. Be-cause the feature size is much larger than the effective wavelength of light, noscattering is expected. For a perpendicular light ray, when it hits the cell, areflected light ray will have a high probability to hit another plane of one ofthe pyramids, which results in a reduction of the reflection from the cell andwill therefore have a higher light absorption. Furthermore, the slanted natureof the pyramids will induce a light path increase, resulting in more absorbedlight. When the angles of the light ray on various interfaces within the cell arelarge enough, the light may be trapped within the active layers of the solar cellthrough total internal reflection. A schematic representation of this schemeis shown in figure 5.4. Using a 3D ray tracing program, explained in the fol-lowing section, we modelled the absorption enhancement of p-i-n silicon layerstacks on such pyramidal structures, while varying the angle of the pyramidplanes with respect to the substrate surface and calculated at which anglesmaximum absorption enhancement takes place. We have done this for bothdirect and diffuse incoming light. For both types of illumination a significantlight absorption enhancement is expected.


Incoming Light Angle

PC

ZnO:Al

a-Si 275 nm

1000 nm

Pyramid angle

Figure 5.4: A schematic 2D representation of the model used for ray tracingcalculations. For the simulations we used a ZnO:Al thickness of 1000 nm andan a-Si thickness of 275 nm. The pyramid angle is defined as the angle betweenthe substrate surface and one of the pyramid faces.

Simulations

Performing optical modelling on features in the (hundreds of) nanometre range,such as textured interfaces, used for scattering, requires complex and time con-suming calculations, such as finite difference time domain calculations [123].Because our pyramids have dimensions much larger than the wavelength oflight we can use ray optics to calculate the absorption of light in the differ-ent layers. Using a model that combines reflection, refraction, and absorption[124], we calculated the absorption enhancement for several periodically tex-tured substrates at all wavelengths of interest. For the wavelength dependentabsorption properties and refractive indices of the silicon and aluminium dopedzinc oxide TCO layers we used data obtained by measuring the optical prop-erties of these thin films deposited for this study and fitting the optical datausing the OJL model [55]. The absorption coefficient is lower for the a-Si de-posited at 130°C compared to the a-Si:H deposited at 180°C that is used for


optimal layer properties, due to the higher band gap of the material depositedat low temperature. In the ray tracing calculations we modelled the solar cellsas follows: (1) A halfspace of air with a refractive index of 1 (2) as a substratewe used PC with a fixed refractive index of 1.5 and no absorption. (3) As aTCO we modelled 1000 nm of ZnO:Al with refractive index and absorptionas shown in figure 4.1. For the amorphous silicon we used the measured op-tical properties of the low temperature (130°C) sample as shown in figure 4.1and a layer thickness (perpendicular to the substrate surface) of 275 nm. Sub-sequently, we changed the shape of the substrate into square based pyramidswith varying pyramid angles. Taking the pyramid angle as the angle betweenthe substrate surface and the pyramid wall, we varied this angle from -60° to60° in steps of 5°, where a negative pyramid angle denotes an inverted pyramidfacing inwards into the substrate. We simulated light hitting the cell in twoways: (1) direct, perpendicular to the substrate surface and (2) diffuse, usinga Lambertian angle distribution. We used light with wavelengths from 400 nmto 800 nm in steps of 10 nm. For every wavelength the following was calcu-lated: (1) The total fraction of light lost through reflection. (2) The fractionof light absorbed by the ZnO:Al. (3) The fraction of light absorbed by thea-Si and (4) the fraction of light lost by transmission through the completecell. For every wavelength we took the average result of 1000 rays hitting thecell on a randomly chosen position on the cell.

Figure 5.5 shows the calculated reflection and absorption by the ZnO:Allayer, the absorption of the a-Si:H layer of a flat cell and of a cell with pyramidangles of 40°, illuminated with direct light. According to the simulations, thepyramidal structure suppresses the reflection over the whole spectral range,which results in an enhanced absorption in the a-Si, which in turn will resultsin a higher generated current. There is also a small increase in absorptionin the ZnO:Al of the front electrode. If we now integrate the absorbed lightover the AM1.5 spectrum we can calculate the expected generated current fordifferent pyramid angles and different types of light (direct or diffuse). Thecalculated current density results for direct light and diffuse light, for pyramidangles from -60° to 60° are shown in figure 5.6. For direct light, the simulatedcurrent density changes from around 9mA/cm2 for a flat surface to a maximumof 13.5mA/cm2 for pyramid angles of around 45° degrees for both positive andnegative pyramid angles. For diffuse light, the current density increases fromaround 8.5mA/cm2 to 12mA/cm2 for pyramid angles of 30° and higher, forboth normal and inverted pyramids. The increase in absorption is mainlycaused by a lower reflection over the whole spectrum, resulting in a relativeincrease in absorption over the whole spectrum and a relatively high absorptionenhancement in the red part (λ > 600 nm) of the spectrum. According to our


4 0 0 5 0 0 6 0 0 7 0 0 8 0 00 . 00 . 20 . 40 . 60 . 81 . 0

Refle

ction

, Abs

orptio

n


Figure 5.5: Simulated total reflection, ZnO:Al absorption and a-Si absorptionin an a-Si:H cell on micropyramid structured substrates with a pyramid angleof 40° and on a flat substrate for direct incoming light.

simulations, for direct light, a current density enhancement of 45% can beachieved at a negative or positive pyramid angle of 45°, compared to a flatcell. For diffuse light the weighted absorption enhancement is 40% at pyramidangles larger than 40°.

SubstratesExperimentally, the pyramids are regularly arranged and have a pyramid angleof 54.7°. Between the pyramids is a 2 µm wide flat surface area, which accountsfor 36% of the sample surface. The production of the mold and the embossingof the PC substrates was carried out by Aquamarijn microfiltration B.V. Themold for embossing was produced from a silicon wafer (100) by applying apositive photoresist and patterning this with a square pattern with a linewidth of 2 µm using UV-lithography. After development of the photoresiststhe wafer was anisotropically etched in a KOH solution at 80°C to create apattern of pyramid structures on the surface. The photoresist was strippedaway and a 200 nm layer of copper was evaporated on the etched wafer andsubsequently 500 µm nickel was electroplated on the wafer (Technic Elevate Ni

5.2. Low temperature solar cells on PC substrates 73

- 6 0 - 4 0 - 2 0 0 2 0 4 0 6 089

1 01 11 21 31 4

D i r e c t L i g h t D i f f u s e L i g h t

Cu

rrent

Dens

ity (m

A/cm2 )

Figure 5.6: Simulated current density of a-Si cells on micropyramid struc-tured substrates for direct light (solid) and diffuse light (dashed) for differentpyramid angles.

5910). The silicon wafer was removed from the nickel by dissolution in KOH.The obtained mold was used for hot embossing of the micropyramids in PC.The hot embossing was done by clamping the mold on a 1 mm thick piece ofPC between two aluminium plates. The clamps were tightened and placed ina vacuum oven. After evacuation the temperature was gradually increased to200°C and kept stable for 1 hour. After that the oven was cooled down to100°C and the substrate was removed from the mold. An SEM image of thestructures is shown in figure 5.3. One type has pyramids facing outward, outof the substrate and the other type of substrate with pyramids facing inwards.

5.2 Low temperature solar cells on PC sub-strates

To be able to deposit thin film cells on plastics, the process temperatureshould never exceed the glass transition temperature of the substrate used,for all the processing steps. For polycarbonate, this transition temperature isaround 145°C [34], which is lower than the experimentally found optimal sub-


strate temperature of 200°C for a-Si deposition [36]. The consequences of lowsubstrate temperature for the electronic quality and optical properties of thea-Si:H material have been discussed in the previous chapter. In this sectionwe will discuss the deposition of front TCOs on flexible substrates and theiradhesion to these substrates. Because we use plastics as a substrate, we willalso check whether there is degassing from the material that could contamin-ate the growing layer or the vacuum system. Because the thermal expansioncoefficient of PC is rather high, compared to the substrate holder material, weemployed a specially designed substrate holder that allows stretching of thesubstrate while it expands, as discussed in chapter 4.

5.2.1 Cells on PC: Experimental issuesTCO adhesion to plastic substrates

Figure 5.7 shows optical microscope images of ZnO:Al deposited directly ontoflat polycarbonate substrates. The ZnO:Al shows cracks. The layers grown at100W RF power in the sputter deposition (right) show less cracks than theones grown at higher powers (300W, left). The sheet resistance of the latterlayers is much higher than is suitable for use as a front TCO in solar cells.Earlier studies of ZnO:Al grown on flexible plastic substrates do not reportadhesion problems [125, 126]. The adhesion difficulties for ZnO on flat PCsubstrates may be attributed to the difference in thermal expansion of ZnO(4×10−6K−1 [127]) and PC (65×10−6K−1[128]). Because the substrate heatsup during the deposition, it expands. After the deposition, when the substratecontracts as it cools down, the ZnO layer cracks under compressive stress. Themicroscope images in figure 5.7 show debris that agree with this hypothesisof compressive stress. Depositions on our micro- and nanopyramid structuredPC substrates do not show adhesion problems, because the textures act asstress relievers. Because glass has a thermal expansion coefficient comparableto that of ZnO (8×10−6K−1), these problems do not occur when depositingZnO on glass substrates.

Plastic substrate degassing

Before a deposition on a substrate can commence, we need to be sure that thereis no degassing from the substrate while heating it up before the deposition orduring the deposition itself. Volatile elements (such as carbon and moisture)from the substrate could contaminate the layer, resulting in material prop-erties different than anticipated. Furthermore, contamination of the reactor


Figure 5.7: Microscope pictures of cracks formed in ZnO:Al layers depositedon flat PC substrates at different powers. The layers grown at 100W (right)show less cracks than the layers grown at 300 W (left).

chamber could have deleterious effects on subsequent depositions. Therefore,we checked whether our PC substrates release gasses when introduced in thevacuum chamber. We did this by monitoring the background pressure overtime after introduction of the substrate into the reaction chamber, while clos-ing the valve to the pumps. Even when there is no substrate present, thebackground pressure will increase over time when the pumps are disconnectedfrom the chamber. Now, if a substrate releases gasses, the increase in pressurewill occur faster. Because every newly introduced substrate (or any object,such as a substrate holder) from outside the vacuum carries moisture with it,the background pressure is expected to rise after introduction into the vacuum,but will quickly drop over time. Figure 5.8 shows the change in backgroundpressure after introduction of a PC substrate into chamber 4 of the ASTERsystem. The heater temperature was set to 180°C, which corresponds to a slowheating of the substrate up to 130°C. Every few minutes we disconnected thepumps from the chamber and monitored the increase of pressure over time,which was higher when a PC sample was inserted than without a sample forthe first 2 hours, but after that both situations show similar behaviour. There-fore we conclude that there is no continuing degassing from the substrates andthat they are therefore safe to use in our deposition system.


1 0 1 0 0 1 0 0 01 E - 6

1 E - 5

1 E - 4

1 E - 3

Pres

sure

(mba

r)

T i m e ( s )

Figure 5.8: Pressure increase over time when the pump valve is closed afterdifferent times of heating of the substrate in vacuum. For comparison thetime dependent pressure increase of an empty chamber is also shown. Thebackground pressure over time (with substrate and substrate holder mounted,in minutes) with the pump valve open shows a usual pressure decrease.


5.2.2 Solar cell results

To test the light trapping abilities of our different structured substrates, wedeposited single junction a-Si solar cells in a p-i-n configuration on the differentsubstrates. As a front TCO we used ZnO:Al, deposited in our RF magnetronsputtering system, SALSA, from a ZnO:0.5%Al2O3 target at room temper-ature. As a BR we used ZnO:Al from a ZnO:2% Al2O3 target. The siliconlayers were deposited in our ultra high vacuum multi-chamber system ASTERby VHF-PECVD at a frequency of 60 MHz, using a showerhead electrode.The p- and n-layers were deposited at 50 MHz in separate chambers. Thesubstrate temperature during deposition was set at 130°C and is expected toincrease 7°C to 8°C during the deposition. The metal contacts were depositedby thermal evaporation of silver and aluminium. The cell structure was asfollows: The substrate, a 1000 nm ZnO:Al front TCO, a double p-structureof nanocrystalline and amorphous silicon of 15 nm (in total), a 275 nm thickintrinsic a-Si layer, a 30 nm thick a-Si n-layer, a 100 nm ZnO:Al BR and asilver/aluminium back contact. After deposition, the cells were annealed in anitrogen environment for 1 hour at 125°C. The area of the cells was 0.16 cm2.

As a reference, cells were deposited on 2 other types of substrates: (1) AsahiU-type natively textured fluorinated tin oxide (SnO2:F) TCO glass and (2)flat glass (Corning Eagle 2000) substrates with an untreated flat ZnO:Al TCOlayer, under the same low temperature deposition conditions. For referencepurposes, a cell deposited on a flat PC surface would be the best candidate,but the stress of the ZnO:Al layers prevented flat PC to be used as a substrate,due to cracking of the ZnO:Al layers, as discussed above.

Figure 5.9 shows the J-V measurements of the cells on flat glass, Asahi-UTCO glass, micropyramid structured PC and on nanopyramid structured PCsubstrates. Table 5.1 shows the electrical properties for all types of cells, asobtained by J-V measurements under dark conditions and under AM1.5 illu-mination after annealing at 125°C for 1 hour in a nitrogen environment. Thecell on micropyramid structured PC shows an increase in Jsc of 22%, com-pared to the cell on a flat glass substrate. The cell on concave nanopyramids(type I) shows a short-circuit current density increase of 22%, whereas theconvex nanopyramids (type II) enhance the current density by 28%. The cur-rent generated by the cells on convex nanopyramid structured PC substratesis slightly higher than that of the reference cells on Asahi U-type TCO. Fig-ure 5.10 shows the spectral response data for three types of cells on PC andthe cells on flat glass and on Asahi U-type TCO glass. All cells on struc-tured or textured substrates show a significantly higher quantum efficiency atwavelengths higher than 400 nm, when compared to the cell on a flat substrate,


- 0 . 5 0 . 0 0 . 5 1 . 0- 1 5

- 1 0

- 5

0

5 f l a t g l a s s m i c r o p y r a m i d n a n o p y r a m i d I n a n o p y r a m i d I I A s a h i U - t y p e T C O

J (mA/c

m2 )

V o l t a g e ( V )

Figure 5.9: Current density - Voltage characteristics for cells deposited ondifferent embossed PC substrates at 130°C. As a reference J-V characteristicsfor cells on flat glass and on Asahi textured TCO glass are also shown.

Substrate type Jsc Voc FF n J0 Rs Rp η

(mA/cm2) (V) (%) (mA/cm2) (Ωcm2) (Ωcm2) (%)

Flat glass 10.14 0.93 64 1.49 5.6×10−10 7.5 1166 5.6Micropyramid 12.35 0.87 59 1.90 4.7×10−8 9.7 801 6.4Nanopyramid I 12.33 0.88 63 1.69 4.7×10−9 9.3 964 6.8Nanopyramid II 12.98 0.89 64 1.98 1.3×10−7 9.0 1116 7.4

Asahi U-type 12.81 0.93 64 1.49 2.0×10−10 9.1 1250 7.6Jsc: short-circuit current density; Voc : open-circuit voltage; FF: fill factor; n: diode qualityfactor; J0: reverse saturation current; Rs: series resistance; Rp: parallel resistance; η: conversionefficiency. Type I: concave pyramids; type II: convex pyramids

Table 5.1: Initial electrical properties of cells deposited at 130°C on flat ZnO:Alcoated glass substrates, micro- and nanopyramid structured PC substrates andAsahi U-type textured glass as measured by J-V measurements under AM1.5illumination and under dark conditions.


4 0 0 5 0 0 6 0 0 7 0 0 8 0 00 . 00 . 20 . 40 . 60 . 81 . 0 f l a t g l a s s

m i c r o p y r a m i d n a n o p y r a m i d I n a n o p y r a m i d I I A s a h i U - t y p e T C O

ECE


Figure 5.10: External collection efficiency measurements of cells deposited ondifferent embossed PC substrates at 130°C. As a reference the measurementsfor cells on flat glass and on Asahi textured TCO glass are also shown.

which can be ascribed to anti-reflective properties and better response in thered part of the spectrum. If we compare the spectral response of the cellson nanopyramids to that of the cell on Asahi TCO, we observe a comparableresponse at wavelengths above 600 nm, a small decrease in quantum efficiencyfor wavelengths between 500 and 600 nm, whereas the response between 400and 500 nm is higher. This results in a total short-circuit current density gen-eration that is slightly higher for the cell on convex pyramids and a bit smallerfor the cell on concave pyramids, compared to the cell on Asahi TCO. Thecell on micropyramid structured PC, when compared to the cell on Asahi U-type TCO-glass, shows a total current density that is about 0.5mA/cm2 lower.This is mainly caused by the difference in response above 500 nm.

In the ultra-violet part of the spectrum, below 400 nm, absorption of lightby the substrate and the TCO dominate spectral response behaviour. The cellon Asahi TCO, which is made of SnO2:F, shows transparency down to 300 nm,whereas ZnO:Al, due to its lower band gap for the samples in this study, cutsoff the light below 350 nm. PC is not transparent for light of wavelengthsbelow 380 nm. This is illustrated in figure 5.11, which shows the transmissionof glass and PC with and without the ZnO:Al front contact and of Asahi U-


3 0 0 4 0 0 5 0 0 6 0 0 7 0 0 8 0 00 . 00 . 20 . 40 . 60 . 81 . 0

Trans

miss

ion


P o l y c a r b o n a t e Z n O : A l o n P C C o r n i n g g l a s s Z n O : A l o n g l a s s A s a h i - U T C O

Figure 5.11: Total transmission of glass and PC substrates with and withoutZnO:Al and the total transmission of Asahi U-type TCO glass.

type TCO glass. These differences in ultra-violet transmission account for thedifference in spectral response in this wavelength region.

The cells on flat glass and on Asahi TCO show a high Voc of 0.93 V. Thisis caused by the high optical band gap of the intrinsic a-Si:H deposited at lowtemperature (1.9 eV), which is attributed to a higher hydrogen content in thefilm. The cell on structured PC shows a lower Voc of 0.87 V, which can beattributed to a higher dark current (J0). The cell on micropyramid structuredPC shows a conversion efficiency of 6.4%, which is around 1% absolute lowerthan the cell on textured Asahi TCO, which has a slightly higher Voc and ahigher fill factor of 64%, compared to 59% for the cell on micropyramid struc-tured PC. The cells on nanopyramid structured concave and convex pyramidsshow an efficiency of 6.8% and 7.4%, respectively, which is slightly lower thanthat of the cell on Asahi TCO, due to a lower Voc.

Our simulation results (figure 5.6) had shown that a current enhancementfor micropyramid structured substrates is possible up to 45%. Experimentally,for a pyramid angle of 54°, which is close to the calculated optimum pyramidangle, a current enhancement of 22% has been observed, compared to cells ona flat glass substrate. The micropyramid structured substrates that we usedhave a flat surface area of around 30%, due to the fabrication method, whichexplains a lower current enhancement than expected for a substrate that iscompletely filled with pyramids. A better coverage, using the same embossing


4 0 0 5 0 0 6 0 0 7 0 0 8 0 00 . 00 . 20 . 40 . 60 . 81 . 0

Refle

ctanc

e


Figure 5.12: Total reflection of a-Si cells deposited on flat glass substrates,on micropyramid structured PC substrates and on both types of nanopyramidstructured PC. All cells were deposited at a substrate temperature of 130°C.

technique, could be achieved by using larger-sized pyramids. According tothe simulations, the increase in absorbed light is mainly caused by a decreasein reflection of light from the top surface of the cells. Figure 5.12 shows themeasured total reflection from the cell on micropyramid structured PC, bothtypes of nanopyramid structured PC and the cell on flat glass. The cell onthe micropyramid structured substrate shows a lower reflection over the wholemeasured spectral range. This results in a generated current density in thecell on structured PC which is comparable to the cell on Asahi TCO, but thecell on structured PC suffers from a lower Voc and a lower FF than the cellon Asahi. The cell on PC has a higher diode quality factor (n) value of 1.90compared to the cell on Asahi (1.49) and has a saturation current densitywhich is 2 orders of magnitude higher than the latter.

Although the silicon layers of the cells are made under identical depos-ition conditions, the qualities of the diodes differ. There are a few possibleexplanations for this. Firstly, defects could be created in the silicon due tothe thermal expansion of the substrate, the coefficient of thermal expansion ofPC, being much higher than that of amorphous silicon. This could result inexternal stress in the silicon layers and thereby induce defects in these layers.Secondly, controlling the substrate temperature is very important for obtain-ing good quality a-Si films. Although good quality films can be deposited even


Defective region

Figure 5.13: Cross-sectional TEM images of a-Si cells deposited on micropyr-amid structured PC substrates. On the left a crack through the completecell is visible. On the right, defective regions are present in the shoulder of apyramid.

at room temperature, tuning of the hydrogen dilution is necessary to secure adevice quality film [129]. When depositing on plastic substrates, intrinsic stressof the layer on the substrate will cause curving of the substrate, resulting in adecrease in heat transfer from the heater to the substrate, which could resultin a lower substrate temperature. Although our depositions were done in aspecially designed substrate holder, which can compensate for the expansionby moving one end of the holder outwards by a pulling spring, when a layeris deposited which shows compressive stress, the bending of the substrate willcause a gap between the substrate holder and the substrate. Thirdly, studieshave shown that the substrate morphology can have an influence on the defectformation during the deposition of a-Si [40]. Defects are formed within theconcave regions of the substrate. Similar defective regions were also found innanocrystalline silicon thin films [130, 131]. Figure 5.13 shows cross-sectionaltransmission electron microscopy (XTEM) images of the cell on micropyramidstructured PC, on which we can identify defective regions: On the left, weobserve a complete crack through the silicon layers. On the right side, nano-cracks (elongated voids) can be identified in the silicon layers near a valley inthe TCO-silicon interface.

5.3. Post-deposition treatments 83

- 0 . 2 0 . 0 0 . 2 0 . 4 0 . 6 0 . 8 1 . 0- 1 5

- 1 0

- 5

0

5

J (mA

/cm2 )

V o l t a g e ( V )

A f t e r s h u n t b u s t i n g I n i t i a l

Figure 5.14: Shunt busting a cell can change a short-circuited (shunted) cell(dashed) to a working cell (solid).

5.3 Post-deposition treatments

5.3.1 Shunt busting

In thin film solar cells, short-circuit paths, also known as shunts, can destroythe diode behaviour of the cell. Especially when cells have very thin siliconlayers or are on very rough surfaces, shunting of cells can drastically bringdown the yield of solar cells. Earlier studies reveal that these shunt paths aremetastable, i.e. shunt paths can be created and removed by applying forwardand reverse bias voltages respectively [132]. The most likely way a shunt pathcan form is through the incursion of Al from the ZnO:Al BR into pinholes ormacroscopic defects, formed due to dust on the surface during the depositionprocess. During shunt-busting, the Al diffuses out of the a-Si [133]. Figure5.14 shows the J-V characteristics of a standard a-Si thin film solar cell beforeand after our shunt-busting procedure, which consists of applying a linearlyincreasing reverse bias voltage from 0V to -5V in 6 seconds. Care is takenso that the PC substrate is not damaged due to overheating during shuntbusting. Although this procedure can recover some of the shunted cells, itdoes not work for all short-circuited cells. For some cells, the PC substrateheats up too much before it can recover, resulting in a damaged substrate anda destroyed cell.


-0.4 -0.2 0.0 0.2 0.4 0.6 0.8 1.0-15

-10

-5

0

5

J (m

A/c

m2)

Voltage (V)

125°C

135°C

140°C

-0.5 0.0 0.5 1.0 1.51E-9

1E-7

1E-5

1E-3

0.1

10

1000

J (m

A/c

m2)

Voltage (V)

125°C

135°C

140°C

Figure 5.15: Dark (left) and light J-V characteristics of a cell on convex nan-opyramid structured PC after annealing at different temperatures. Annealingup to 135°C improves the quality of the solar cell, whereas after annealing at140°C the cell performance collapses. The slight shift of the minimum fromzero voltage for the curves under dark conditions is an experimental artefact.

5.3.2 Post deposition annealingFigure 5.15 shows the J-V-characteristics of cells deposited on convex nanopyr-amid structured PC, after annealing at different temperatures. Directly afterthe deposition, annealing was done in a nitrogen environment for 4 hour at125°C. Subsequently, the cell was annealed at 135°C for 4 hours. The third andlast annealing step was done at 140°C for 90 minutes. Annealing at 125°C and135°C increases the performance of the cell, which is reflected in the fill factor.After the annealing at 140°C, the fill factor collapses and the cell performancedrops drastically. Although this is below the glass transition temperature ofthe PC material of 145°C [34], the cell is still adversely affected. This maybe due to structural changes of the PC below the glass transition temperature[134].

5.3.3 Stability under light soakingDegradation of a-Si solar cells under light soaking conditions reduces the cellperformance over time [25]. To test the light induced degradation of a certaincell, a standard test is undertaken in which the cell is soaked under AM1.5-like illumination conditions for 1000 hours, while keeping the temperatureconstant at 50°C under open-circuit conditions. As the metal-halide lamp

5.4. Conclusions 85

- 0 . 4 - 0 . 2 0 . 0 0 . 2 0 . 4 0 . 6 0 . 8 1 . 0- 1 5

- 1 0

- 5

0

5

J (mA

/cm2 )

V o l t a g e ( V )

Figure 5.16: J-V characteristics of an a-Si cell on nanopyramid structured PCbefore and after light soaking under AM1.5-like illumination conditions for1000 hours. A re-annealing step improves the Voc and FF.

degrades gradually while in use, the intensity (in this case) was roughly 70%the intensity of the AM1.5 spectrum. Normally, we would monitor the cellduring the complete soaking time, but as the cells on PC are easily damagedduring the measurements, we chose to only measure the cell performance beforeand after light soaking. Figure 5.16 shows these measurements. After lightsoaking, the Voc drops from 0.88V to 0.85V and the fill factor drops from0.66 to 0.56. After a re-annealing step at 125°C for 1 hour, the Voc increasesto 0.88V and the fill factor increases up to 0.61.

5.4 ConclusionsWe studied the light trapping in a-Si cells deposited on PC substrates usingthree different substrate structures: Asahi U-type, nanopyramids and pyram-ids much larger than the wavelength of light. We have achieved an initial solarcell efficiency of 7.4% for a cell directly deposited on nanopyramid structuredPC after annealing for 1 hour at 125°C. Compared to cells on Asahi U-type,these cells show a slightly higher current density, but suffer from a lower Voc.Cells on micropyramid structured PC had a maximum initial efficiency of 6.4%after annealing, which is lower due to a lower current and a lower fill factor.


XTEM studies show that the cells deposited on PC substrates have more de-fects than the cells grown on glass substrates, which could be caused by a largedifference in thermal expansion of PC and a-Si, or by the sharp features of thestructured PC samples. This is reflected in a higher reverse saturation currentfor cells deposited on PC.

Post-annealing improves the efficiency of a-Si cells deposited at low temper-ature. Some short-circuited cells can be recovered by applying a shunt-bustingprocedure, in which a reverse bias is applied for several seconds.

Chapter 6

Micromorph tandem cellson plastic substrates

6.1 IntroductionNanocrystalline silicon (nc-Si:H) differs from amorphous silicon (a-Si:H) in anumber of ways. A very important property for solar cell applications is itslower band gap (1.1 eV) compared to a-Si:H (1.8 eV), enabling it to absorb lightwith wavelengths up to ∼1100 nm. A stacked structure of an a-Si:H cell anda nc-Si:H cell is called the micromorph concept, which was first introduced byIMT (now EPFL) [13]. This concept allows the blue and green part of thespectrum to be absorbed by the a-Si:H, the top cell, whereas the remaininglight is passed on to the nc-Si:H cell, the bottom cell, which will absorb mostlyred light. In this way, the mismatch between the energy of the absorbedphotons and the band gap is minimized, and therefore thermalization lossesare reduced, compared to a nc-Si:H cell. Compared to a single junction a-Si:Hcell, light of a broader spectrum is absorbed.

The absorption coefficient of nc-Si:H is low, compared to amorphous mater-ial with a similar band gap, such as a-(Si)Ge, due to a predominantly indirectnature of the band-gap. Therefore rather thick absorber layers are used forsolar cells, with typical thicknesses of 1 to 3µm. Because in a micromorphtandem cell the individual cells are connected in series, the total cell current isequal to the lowest current generated by one of the two cells. Therefore goodcurrent matching between the sub-cells in a tandem cell is crucial for goodelectrical performance.

88 Chapter 6. Micromorph tandem cells on plastic substrates

In this chapter, we report on our studies towards depositing nc-Si:H and a-Si:H/nc-Si:H tandem cells at a substrate temperature of 130°C, and our resultson direct depositions of cells onto polycarbonate (PC) substrates. For this weneeded to adapt the deposition processes to obtain device quality materialsat low temperature. The material studies on the layers that are used in thischapter are described in chapter 4. For the a-Si:H top cells, the cells treatedin chapter 5 are used as a basis.

As described in chapter 4, to obtain device quality material at low tem-peratures, silicon layers have to be deposited at a higher hydrogen to silanegas flow ratio (hydrogen dilution) than that is required at higher substratetemperatures [37]. As a result, deposition rates decrease, resulting in longerdeposition times. This time factor becomes even more severe when we considermulti-junction cells, because of the thick hydrogenated nanocrystalline siliconbottom cell needed to achieve adequate current matching between the top andthe bottom cell. We tackle this problem by reducing the total thickness of thecell to around 1000 nm. This concept has been applied to cells at high temper-atures [135], sometimes making use of intermediate reflecting layers betweenthe top and bottom cells [136].

6.2 nc-Si:H cells on glass substratesBased on the nc-Si:H intrinsic layer series as a function of applied plasmapower input, as described in chapter 4, we produced solar cells on glass in ap-i-n configuration. As a front transparant conducting oxide (TCO) we usedaluminium doped zinc oxide (ZnO:Al) which was sputter-deposited, followedby texture etching in a hydrochloric acid solution. As a back contact weused evaporated silver, after sputter-depositing a ZnO:Al back reflector. Thethickness of the intrinsic layers of the cells was aimed at 700 nm, based ondeposition rate. Before characterization, the cells were annealed in a nitrogenenvironment for 1 hour at 125°C. The size of the cells was 0.16 cm2.

Figure 6.1 shows the current density-voltage (J-V) characteristics of nc-Si:H cells, as a function of applied plasma power around the transition froma-Si:H to nc-Si:H. Going from 12.5W to 22.5W, the material (as shown infigure 4.8) changes from mostly amorphous to almost fully crystalline. Figure6.2 shows the short-circuit density (Jsc), open-circuit voltage (Voc), fill factor(FF) and the resulting conversion efficiency (η) before and after annealing for1 hour at 125°C. For the cells made from the layers in this series this results inan increasing current density up to 20W applied plasma power, which can beattributed to a lower band gap for the material made at at higher power, owing

6.2. nc-Si:H cells on glass substrates 89

- 0 . 4 - 0 . 2 0 . 0 0 . 2 0 . 4 0 . 6 0 . 8- 2 0- 1 5- 1 0

- 505

Curre

nt De

nsity

(mA/c

m2 )

V o l t a g e ( V )

1 2 . 5 W 1 5 W 1 7 . 5 W 2 0 W 2 2 . 5 W

Figure 6.1: J-V curves for nc-Si:H cells deposited at 130°C on glass substratesas a function of applied VHF power. The structure of the cell is glass/20 nmp-nc-Si:H/700 nm i-nc-Si:H/30 nm n-a-Si:H/100 nm ZnO:Al/Ag/Al.

to a larger crystalline fraction. At 20W power, the i-layer is fully crystallineand the band gap for the material made at higher powers does not decreaseany more. At the same time, the Voc decreases from 0.62V to 0.49V for theannealed cells. The FF shows a maximum at 17.5W applied power, where thecrystalline fraction is roughly 40%. The resulting efficiencies show the sametrend, peaking at 17.5W. At this point, annealing raises the Voc by about60meV, while the FF is lifted by about 8% absolute. The highest achievedconversion efficiency is 6.9%.

As mentioned in the introduction, deposition rate becomes an issue at lowdeposition temperature. We have been successful in depositing nc-Si:H ata reasonably high deposition rate (0.51 nm/s) using high pressure (3 mbar)and high power and a showerhead cathode for the gas distribution in theplasma zone for the deposition. This is an adaptation of the deposition processwhich delivered 10% efficiency nc-Si:H cells at standard deposition temperature(180°C) [137].


1 0 1 5 2 0 2 51 2

1 4

1 6

1 8

2 0

2 2

1 0 1 5 2 0 2 5 0 . 4 0

0 . 4 8

0 . 5 6

0 . 6 4

1 0 1 5 2 0 2 5

0 . 5 5

0 . 6 0

0 . 6 5

1 0 1 5 2 0 2 54

5

6

7

Curre

nt De

nsity

(mA/

cm2 )

Open

Circ

uit Vo

ltage

(V)

Fill F

actor

(%)

P o w e r ( W ) Ef

ficien

cy (%

)

P o w e r ( W )

Figure 6.2: Electrical properties of nc-Si:H cells deposited at 130°C on glasssubstrates as a function of applied VHF power, before and after an annealingstep at 125°C of 1 hour.

6.3. Tandem cells on glass substrates 91

4 0 0 5 0 0 6 0 0 7 0 0 8 0 0 9 0 0 1 0 0 00 . 00 . 20 . 40 . 60 . 81 . 0

ECE


7 0 0 n m B o t t o m C e l l 9 0 0 n m B o t t o m C e l l

T o p C e l l s B o t t o m C e l l s

T a n d e m T o t a l s

Figure 6.3: Spectral response characteristics for tandem cells deposited on glasssubstrates at 130°C, showing cells with two different bottom cell thicknesses:700 nm (solid) and 900 nm (dashed). The spectral response for the top cell,bottom cell and summed up response are shown individually.

6.3 Tandem cells on glass substrates

The single junction a-Si:H cells as described in chapter 4 together with the nc-Si:H cells from the previous section were combined to produce a-Si:H/nc-Si:Htandem cells at 130°C. A double p-layer (nc-Si:H/a-Si:H) was used to makeproper contact with the texture etched ZnO:Al front TCO. The complete solarcell structure was as follows: Superstrate/ZnO:Al TCO/p nc-Si:H/p a-Si:H/ia-Si:H/n a-Si:H/p nc-Si:H/i nc-Si:H/n a Si:H/ZnO:Al/Ag/Al.

We used very thin layers as the photo-active layers: 275 nm of a-Si:H forthe top cell, combined with a nc-Si:H bottom cell with an i-layer of 700 nm.Figure 6.3 (solid) shows the external collection efficiency (ECE) data for theresulting tandem solar cell structure, showing the spectral response for thetop a-Si:H cell and the bottom nc-Si:H cell separately and the total (sum)ECE of the structure. The calculated current densities for the cells showeda mismatch between the top and bottom cells: 8.8mA/cm2 for the top celland 7.6mA/cm2 for the bottom cell. Based on these measurements we rede-posited the solar cell structure, but now using a thicker bottom cell of 900 nm


- 0 . 5 0 . 0 0 . 5 1 . 0 1 . 5- 2 0- 1 5- 1 0

- 505

a - S i , 2 7 5 n m n c - S i , 7 0 0 n m t a n d e m , B C 7 0 0 n m t a n d e m , B C 9 0 0 n m

J (mA

/cm2 )

V o l t a g e ( V )

Figure 6.4: J-V characteristics for cells deposited on glass substrates at 130°Cunder AM1.5 light conditions, showing cells with two different bottom cellthicknesses: 700 nm (dashed) and 900 nm (solid). Also the J-V characteristicsof the single junction cells on which the tandem was based are shown.

for better current matching. The results are also in figure 6.3 (dashed). Cal-culated from the spectral response measurements, the bottom cell currentincreased to 9.1mA/cm2, whereas the top cell current also increased slightly,to 9.2mA/cm2. The increase in top cell current is probably due to a slightchange in deposition conditions.

Figure 6.4 shows the J-V curves under AM1.5 illumination of the a-Si:H cell(275 nm) and the nc-Si:H cell (700 nm) on which the deposition recipe for thetandem was based, and the resulting tandem cell. Also shown is the tandemcell with an i-layer thickness of 900 nm for the bottom cell. The J-V data forthe a-Si:H cell shows a rather high Voc of 0.90V, which is attributed to thehigh band gap of the a-Si:H layer deposited at 130°C. The characteristics of thecurves, especially near Voc, confirms that the tunnel recombination junctionwith the low temperature doped layers is working well, showing no S-character.The micromorph tandem cell has a Voc of 1.40V, a Jsc of 10.5mA/cm2 and aFF of 65%, resulting in an initial efficiency of 9.5%.

Figure 6.5 shows a bright field cross-sectional transmission electron mi-croscopy (XTEM) image of the complete cell structure showing the different

6.3. Tandem cells on glass substrates 93

layers. Analysis of the XTEM images shows a top cell thickness of 285± 10nm and a bottom cell thickness of 960± 60 nm, including p- and n-layers. Thedeposition rates of the a-Si:H and the nc-Si:H i-layer are 0.22 nm/s and 0.51nm/s, respectively, which results in a total deposition time for the i-layers ofonly 51 minutes. The deposition time for all Si layers of the tandem cell is justover 80 minutes. The p- and n-layers were not optimized for deposition speed.Apart from the greatly reduced deposition time, reducing the thickness of thebottom cell has a number of advantages. First of all, it results in a decreasein material usage. Secondly, the thinner layers mitigate the deleterious effectof a relatively high defect density (resulting from deposition at lower than op-timum temperature) on FF and Voc. Thirdly, thinner layers induce less stresson the substrate, which is a very important property when using plastics as asubstrate.

Cost reduction (together with flexibility) is the main advantage of usingplastics as a substrate material. Apart from the substrate, cost reductionscan be achieved by speeding up the manufacturing process (througput), i.e.reducing deposition times. This can be achieved by either increasing the de-position rate of the layers or by decreasing the thickness of the layers. Inthis study we have shown that excellent spectral splitting can be achieved invery thin micromorph tandem solar cells deposited at 130°. Because the lowtemperature a-Si:H has a high band gap of 1.9 eV, corresponding to light witha wavelength of 653 nm, the top cell will transmit more (red) light towardsthe bottom cell than an a-Si:H cell deposited at higher temperature, whilethe Voc of the completed cell will rise. This results in the possibility to use athinner bottom cell, both because there is more light available for the bottomcell and because of the lower current generated by the top cell. The resultinglower Jsc of the complete cell is partly compensated by the higher Voc. In ourconfiguration, we achieved excellent current matching using a 275 nm thicka-Si:H top cell and a 900 nm thick nc-Si:H bottom cell.

6.3.1 Stability under light soakingDegradation of solar cells under light conditions reduces the cell performanceover time [25]. This is especially true for a-Si:H solar cells. To test the lightinduced degradation of a low temperature tandem cell, a standard test is un-dertaken in which the cell is illuminated by an AM1.5-like spectrum, in ourlight soaking set up, which is described in Chapter 2. The power densityapproaches 100mW/cm2. The temperature of the samples is controlled andkept at a constant 50°C. During the light soaking, J-V measurements are per-formed at exponentially increasing time intervals. Between the measurements,


nc-Si:H bottom cell

a-Si:H top cell TRJ

ZnO:Al

ZnO:Al Al/Ag

Figure 6.5: XTEM image of the a-Si:H/nc-Si:H tandem solar cell, showing theZnO:Al TCO, a-Si:H top cell and nc-Si:H bottom cell. A clear boundary isvisible where the tunnel recombination junction (TRJ) is between the a-Si:Htop cell and the nc-Si:H bottom cell. The cells have total thicknesses of 285 nmfor the top cell and 960 nm for the bottom cell.

6.4. Tandem cells on plastic substrates 95

0 . 0 1 0 . 1 1 1 0 1 0 0 1 0 0 00 . 9 00 . 9 20 . 9 40 . 9 60 . 9 81 . 0 0

No

rmali

zed V

oc, F

ill Fa

ctor

L i g h t S o a k i n g T i m e ( H o u r s )

V o c F F

Figure 6.6: The normalized degradation of the FF and Voc of an a-Si:H/nc-Si:H tandem cell deposited at 130°C.

the cells are kept under open-circuit conditions. The light intensity during thelight soaking was monitored using a crystalline silicon reference diode.

Figure 6.6 shows these measurements for a-Si:H/nc-Si:H tandem cells de-posited at a substrate temperature of 130°C on glass. We can see that the Vocdecreases roughly 7.5% and the FF decreases 9% over time. The main partof the degradation of the fill factor occurs within 10 to 100 hours, whereasof the Voc degradation occurs within the first hours of light soaking. Thefast degradation of the Voc implies that a part of the degradation has alreadytaken place when the J-V characteristics of the cells were measured in oursolar simulator. The different degradation times indicate different degrada-tion mechanisms. The degradation of the fill factor is related to the formationof dangling bonds, which act as recombination centres, which cause a decreasein current, especially under forward bias-conditions.

6.4 Tandem cells on plastic substratesFor the deposition of an a-Si:H/nc-Si:H tandem cells on polycarbonate wecopied the recipe from the previous section and performed a deposition run


- 0 . 5 0 . 0 0 . 5 1 . 0 1 . 5- 1 5 . 0- 1 2 . 5- 1 0 . 0

- 7 . 5- 5 . 0- 2 . 50 . 02 . 55 . 07 . 5

4 0 0 6 0 0 8 0 0 1 0 0 00 . 00 . 20 . 40 . 60 . 81 . 0

J (mA

/cm2 )

V o l t a g e ( V )

Φ S i H 4/ Φ H 2

= 5 / 1 0 0 Φ S i H 4

/ Φ H 2 = 4 . 5 / 1 0 0

ECE


N o B i a s L i g h t B o t t o m c e l l T o p c e l l

Φ S i H 4 / Φ H 2 = 4 . 5 / 1 0 0

Figure 6.7: (left) J-V curves under AM1.5 light conditions of tandem cellsdeposited on micro-structured PC substrates, using two different hydrogendilutions for the bottom cell i-layer. The inset shows Raman spectroscopydata on two bottom cell i-layers at ΦSiH4/ΦH2 =5/100 on glass and on PC.(right) Spectral response data for the tandem cell deposited at the higherhydrogen dilution, showing the response of the top cell, the bottom cell andthe response under dark conditions.

using different types of structured PC, described in chapter 5, as substrates inthe stretch substrate holder. Figure 6.7 (left, solid) shows the J-V measure-ments under AM1.5 light conditions on the resulting cell on micro-structuredPC. The high Voc (1.58V) and the very low current density (3.7mA/cm2)point towards an a-Si:H bottom cell, whereas for the same deposition run on aglass substrate, the bottom cell showed nc-Si:H behaviour. The a-Si:H growthon PC substrates was confirmed by Raman spectroscopy measurements, shownin the inset, which indeed show only a very low Raman crystalline ratio (0.07),whereas the cell deposited on a glass substrate shows a Raman crystalline ratioof 0.47. For this purpose we deposited Si:H layers using the same recipe as thebottom cell i-layer on top of a ZnO:Al/p nc-Si:H structure on micro-structuredPC.

Consequently we deposited a tandem cell structure on different types ofstructured PC, but now changing the gas flows from ΦSiH4/ΦH2 =5/100 toΦSiH4/ΦH2 =4.5/100. Figure 6.7 (left, dashed) shows the J-V curves of theresulting cell on micro-structured PC, showing a higher Jsc of 7.9 mA/cm2 anda Voc of 1.25 V, indicating that now the bottom cell is indeed nanocrystalline.

6.4. Tandem cells on plastic substrates 97

- 0 . 5 0 . 0 0 . 5 1 . 0 1 . 5- 1 0

- 5

0

5

J (mA

/cm2 )

V o l t a g e ( V )

G l a s s , d B C = 9 0 0 n m P C , d B C = 9 0 0 n m P C , d B C = 1 3 5 0 n m

Figure 6.8: J-V measurements of tandem cells deposited on micro-structuredPC, using two different bottom cell thicknesses. As a comparison, the data fora low temperature cell on glass (900 nm bottom cell) in also shown. The insetshows the corresponding dark J-V data.

Figure 6.7 (right) shows the spectral response curves for the same cell. Thegenerated current in the bottom cell is much lower (6.9mA/cm2) than thecurrent from the top cell (8.2mA/cm2) and therefore the current in the cellis limited by the current generated in the bottom cell. For the cells depositedon both types of nanopyramid structured PC (as described in chapter 5), thecurrent generated in the bottom cells is lower, between 5.5 and 6mA/cm2,because the nanostructures are designed for light trapping in a-Si:H cells. Forlight trapping in the red and infra-red part of the spectrum, larger sized pyr-amids are needed [112]. Coming back to our working a-Si/nc-Si tandem cell onmicro-structured PC, the black line in figure 6.7 shows the spectral responseunder dark conditions. The measured spectral response in the blue light regionunder dark conditions, which does not follow the spectral response of the bot-tom cell in this region, indicates that the bottom cell can conduct current evenwhen it is not illuminated. When this happens, we say that the cell is leaking[138], which is probably due to low-quality nc-Si:H which can have shunt pathsand/or a high midgap defect density in the layer, which has adverse effects onthe Voc and fill factor of the cell [139].


4 0 0 5 0 0 6 0 0 7 0 0 8 0 0 9 0 0 1 0 0 00 . 00 . 20 . 40 . 60 . 81 . 0

C e l l T o t a l s

B o t t o m C e l l s

ECE


T a n d e m o n g l a s s , d B C = 9 0 0 n m T a n d e m o n P C , d B C = 9 0 0 n m T a n d e m o n P C , d B C = 1 3 5 0 n m

T o p C e l l s

Figure 6.9: Spectral response measurements on tandem cells deposited on PC,using two different bottom cell thicknesses. As a comparison, the data for thebest low temperature tandem cell on glass in also shown.

To obtain a better matching of the currents generated by the top andbottom cell we deposited new tandem cells, but now using a bottom cell i-layerthickness of 1350 nm, as opposed to the 900 nm used in the previous runs. It isalso believed that a thicker bottom cell will reduce the leaking of the bottomcell. Figure 6.8 shows the J-V characteristics under AM1.5 illumination forthree a-Si:H/nc-Si:H tandem cells: two cells deposited on micro-structuredPC substrates, utilizing two different bottom cell thicknesses, and the tandemdeposited on glass substrates (900 nm bottom cell), as described in the previoussection. Figure 6.9 shows the corresponding spectral responses. Although theresponse of the bottom cell increases (8.3mA/cm2) when a thicker bottomcell is used, the current of the complete cell is now limited by the top cell(7.2mA/cm2). The lower current of the top cell is caused by a lower spectralresponse in the 350-500 nm region. This is probably due to a thicker p-layeror a less transparent TCO layer. Unfortunately, unforeseen circumstancesprevented us from redepositing this run.

When we compare the cells deposited on PC to the cell on glass, we observea lower current generated in the bottom cell. The low voltages for tandem cellson PC can be attributed to the leakage in the bottom cells. As opposed to

6.5. Conclusions 99

Substrate type dBC η Jsc Voc FF Rs Rp JT C JBC

nm (%) (mA/cm2) (V) (%) (Ωcm2) (Ωcm2) (mA/cm2) (mA/cm2)

Glass/TE ZnO:Al 700 9.0 9.7 1.37 68 14.9 2800 8.6 7.6

Glass/TE ZnO:Al 900 9.5 10.5 1.40 65 15.9 2008 9.2 9.1

Micro-struct. PC 900 6.1 7.9 1.25 61 20.4 2172 8.2 6.9

Micro-struct. PC 1350 5.3 6.5 1.26 65 23.0 3426 7.2 8.3

Nano-struct. PC I 900 5.6 6.2 1.27 72 21.0 7025 8.0 5.5

Nano-struct. PC I 1350 5.6 8.0 1.19 59 20.7 1449 n/a 8.3

Nano-struct. PC II 1350 5.2 6.4 1.27 65 41.0 6925 7.9 5.8dBC : bottom cell thickness; η: conversion efficiency; Jsc: current density; Voc: open-circuitvoltage; FF: fill factor; Rs: Series resistance; Rp: parallel resistance; JT C : top cell currentdensity; JBC : bottom cell current density

Table 6.1: Electrical properties of a-Si:H/nc-Si-H tandem cells deposited onglass and on micro- and nano-structured PC at a substrate temperature of130°C. Values are obtained from J-V measurements under AM1.5 illuminationand from spectral response measurements. All top cell i-layers are 275 nmthick.

the a-Si:H cells described in chapter 4, the micro-structured substrates do notexhibit light trapping comparable to texture-etched ZnO:Al in nc-Si:H cells.The cells on micro-structured PC show a decrease in Voc, compared to the cellson glass. When we look at the J-V curves measured under dark conditions(figure 6.8, inset), we see that the cell on glass has a lower diode qualityfactor n of 3.5, compared to the cells on PC (∼3.9) and a reverse saturationcurrent of 9.1× 10−10, which is about one order of magnitude lower than thecells deposited on PC. This indicates a lower material quality of the materialsdeposited on PC. A similar observation was made for a-Si:H single junctionsolar cells deposited on PC substrates.

An overview of the solar cell properties of all tandem cells described in thischapter is given in table 6.1.

6.5 ConclusionsIn this chapter we presented nc-Si:H and a-Si:H/nc-Si:H tandem solar cellsdeposited on glass and on PC substrates. We optimized the nc-Si:H layerquality by tuning the hydrogen dilution when depositing the layer and wereable to accurately control the Raman crystallinity ratio by changing the ap-plied power into the plasma. These layers were used to fabricate nc-Si:H cells


on glass substrates, which gave us good cell performance, using thin (700 nm)i-layers. Combined in an a-Si:H/nc-Si:H tandem cell structure, we were ableto deposit a tandem solar cell at a substrate temperature of 130°C with an(initial) conversion efficiency of 9.5%. Because the used silicon layers in thiscell are rather thin, deposition times can be kept to a minimum, which resultsin a total deposition time for all silicon layers of 80 minutes, and less than anhour for the i-layers only.

When this recipe was transferred to PC substrates, the crystal nucleationbehaved differently, resulting in a-Si:H/a-Si:H tandem cells, which suffer froma very low current, due to bad current matching. Changing the hydrogendilution solved this problem. The light trapping in structured PC substratesis less pronounced in the bottom cell, compared to texture etched ZnO:Al,which is used for light trapping in our tandem cells on glass. Therefore wedeposited tandem cells with a thicker bottom cell to obtain better currentmatching.

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Summary

In the search for new and renewable energy sources, solar energy can fulfila large part of the growing demand. The biggest threshold for large-scalesolar energy harvesting is the price of the solar panels, of which at presentthe workhorse is the crystalline silicon solar cell, made from silicon wafers. Amethod to decrease solar panel prices is the use of thin-film solar cells, whichrequire only a fraction of the raw material compared to crystalline silicon cells.For further cost reductions, continuous fabrication using low-cost substratescan be a solution. In this thesis, we investigate the possibilities of depositingthin film film solar cells directly onto cheap plastic substrates. Apart from thelow cost, flexible solar cell can be used on speciality products, such as clothingor security papers.

Thin silicon films are deposited from the gas phase, in our case using (veryhigh frequency) plasma-enhanced chemical vapour deposition. In this process,feedstock gasses (silane and hydrogen) are decomposed in a reaction chamberusing a radio frequency discharge. Decomposition of these gasses producesradicals, which can reach the substrate, where a thin film grows. In these typesof processes, the temperature of the growing surface has a large influence onthe quality of the grown films.

Because plastic substrates limit the maximum tolerable substrate temper-ature, new methods have to be developed to produce device-grade silicon lay-ers. But lowering the substrate temperature does not only alter the behaviourof depositing species on the growth surface of the films grown, it also changesthe behaviour of the plasma inside the reaction chamber. Apart from growinga film on the substrate, silyl radicals can also polymerize in the plasma bulk.If the resulting particles are negatively charged, they are trapped inside theplasma and continue to grow. At a certain size and density, they will coagu-late and form dust particles in the plasma, which can cause a serious threat todevice operation if they are captured in the thin films. We have shown that us-ing an all-optical technique, we can identify whether dust particles are present

116 Summary

in the plasma or not. For this, we study the axial optical emission from theplasma (caused by relaxation of excited state of species in the plasma). Dueto the presence of the particles in the plasma that capture electrons, the elec-tron temperature increases, inducing a higher optical emission. Because the(large) dust particles encounter gravity, they are pulled towards the bottomof the reactor and therefore the plasma shows an asymmetric optical emissionprofile.

To understand why the formation of dust is temperature-sensitive, we mon-itored the formation of polysilanes as a function of temperature using a mass-spectrometer attached to a plasma reactor chamber. Counting the differentpolysilane radicals at different temperatures in a dust-free regime, but in thedust forming incubation phase, we found the polymerization rate to be in-fluenced by the substrate temperature, which can explain the temperaturedependence of dust formation.

As a substrate material for solar cells, we chose polycarbonate (PC), be-cause of its excellent transparency and its relatively high glass transition tem-perature of 130-140°C. At 130° we searched for deposition recipes that yieldgood quality silicon layers. For this purpose we first investigated how we canaccurately control the substrate temperature. Diluting the feedstock silane gaswith hydrogen has a large influence on the material properties. In the case ofamorphous silicon (a-Si), increasing the hydrogen dilution generally improvesthe quality of the silicon until we reach the nanocrystalline silicon (nc-Si) re-gime. Just before this regime, the a-Si layers show high intrinsic stress, whichmight result in detachment of the silicon layer from the substrate. In the nc-Siregime, together with changing the power input into the plasma, the hydrogendilution can be used to control the volume fraction of crystallites within thesilicon layer.

Using these silicon layers, including doped silicon layers at low temperature,a-Si thin film solar cells were fabricated with an intrinsic layer thickness of275 nm, both on glass and PC substrates. Because low temperature silicon isgenerally not as good as its high temperature counterpart, recombination ofphotogenerated charge carriers can be a problem, resulting in a lower Voc andfill factor. These problems can be mitigated when thinner silicon layers areused and therefore an adequate light trapping technique needs to be employed.For a-Si cells, we have simulated and experimentally tested three light trappingtechniques, using embossed structures in PC substrates and random structureson glass, using features of different sizes: regular pyramid structures largerthan the wavelength of light (micropyramids), pyramid structures comparableto the wavelength of light (nanopyramids) and random nano-textures as usedin commercial fabrication (Asahi U-type TCO glass). Both micropyramid

Summary 117

and nanopyramid substrates enhance the light absorption within the cells.Using nanopyramid substrates we could achieve current densities in cells onPC comparable to current densities achieved on Asahi U-type TCO glass.Using these techniques we could achieve initial conversion efficiencies for a-Sicells on PC of 6.4% on micropyramid substrates and 7.4% on nanopyramidsubstrates, compared to 7.6% for cells deposited under identical conditions onAsahi U-type TCO glass. For nc-Si cells on texture etched aluminium dopedzinc oxide (ZnO:Al) on glass, we could achieve an initial conversion efficiencyof 6.9% using a very thin absorber layer of 750 nm.

Combining low temperature a-Si and nc-Si cells we fabricated tandem solarcells in the ’micromorph’ concept at 130°C. By optimizing the thicknesses ofthe different silicon layers and controlling the crystalline fraction of the bottom(nc-Si) cell, we could achieve an initial conversion efficiency of 9.5% on textureetched ZnO:Al coated glass. When this recipe was transferred to structuredPC substrates, the crystal nucleation behaved differently, resulting in a-Si/a-Si tandem cells, which suffer from a very low current, due to bad currentmatching. Changing the hydrogen dilution could solve this problem. Thelight trapping in structured PC substrates is less pronounced in the bottomcell, compared to texture etched ZnO:Al, which is used for light trapping inour tandem cells on glass.

Samenvatting

Op zoek naar nieuwe betaalbare alternatieve energiebronnen is zonne-energieeen veelbelovende. De aarde ontvangt meer dan 1000 keer zoveel energie daner verbruikt wordt. De grootste belemmering voor grootschalige oogst vanzonne-energie is de prijs van zonnepanelen, waarvan het meestgebruikte typede zogenaamde kristallijne zonnecel is, gemaakt van in plakken gezaagde brok-ken extreem zuiver en daardoor duur silicium. Een alternatief voor deze tech-niek is het gebruik van dunne films, waarvan de productie slechts een frac-tie van het materiaal van zijn kristallijne tegenhanger nodig heeft. Verderekostenbesparingen zouden kunnen worden geboekt door goedkope substratente gebruiken, het liefst in een continu proces. Dit proefschrift beschrijft mijnonderzoek naar technieken om dunne-film silicium zonnecellen direct te produ-ceren op goedkope flexibele plastic substraten. Los van de kostenbesparingen,zouden flexibele zonnecellen niche-markten kunnen bedienen, zoals zonnecellenop kleding of waardepapieren.

Dunne silicium films worden gefabriceerd in vacuümreactors, waar eenplasma de procesgassen silaan (SiH4) en waterstof (H2) ontleedt tot radi-cale moleculen, die wanneer ze op het substraat neerslaan, daar langzaameen dunne laag vormen. Deze radicalen kunnen, buiten een laag vormen, ookaan elkaar plakken, zodat in het plasma materieklontjes of ’stof’ ontstaat. Ditproces van stofvorming blijkt erg temperatuurgevoelig te zijn. Omdat ons uit-eindelijk doel het fabriceren van zonnecellen op plastic substraten is, kunnenwe niet met hoge temperaturen werken. Omdat we hebben gekozen voor po-lycarbonaat (PC) (omdat het hoogtransparant is en verwerkt kan worden bijrelatief hoge temperaturen) mag de temperatuur niet hoger worden dan 130°Com de substraten niet te beschadigen. De eerste stap bij het ontstaan van stofis het groeien van silaanclusters. Als deze clusters negatief geladen zijn, wordende deeltjes ’gevangen’ in de depositiereactor, waar ze verder kunnen groeien.Bij een bepaalde grootte en dichtheid van de clusters klonteren ze samen envormen stofdeeltjes, die de werking van de gedeponeerde zonnecellen kunnen

120 Samenvatting

saboteren. We laten zien dat we kunnen aantonen of er stofdeeltjes aanwezigzijn in het plasma, op basis van een optische techniek. Hiervoor bestuderenwe de optische emissie (door relaxatie van aangeslagen atomen en moleculen)van het plasma als functie van de positie in het plasma. Stofdeeltjes vangenelektronen in, waardoor de elektrontemperatuur lokaal stijgt, resulterend ineen hogere optische emissie. Doordat de stofdeeltjes zich naar de onderkantvan de reactor begeven door de zwaartekracht, verraden zij hun aanwezigheiddoor een asymmetrisch uitgezonden emissieprofiel.

Met een massaspectrometer bestudeerden we de grootte en en aanwezig-heid van clusters in het plasma, als functie van de temperatuur, om zo inzichtte krijgen in de temperatuurafhankelijkheid van stofvorming. Er bleek eentemperatuurafhankelijkheid te bestaan van de groei van de clusters, wat kanverklaren waarom plasma’s bij lage temperatuur eerder stofproducerend wor-den.

Bij een substraattemperatuur van 130°C (de maximaal toelaatbare tem-peratuur voor het gebruik van PC) zochten we naar depositiemethodes voorsiliciumlagen van goede kwaliteit. Hiervoor was het belangrijk de substraat-temperatuur nauwkeurig te kunnen beheersen. In het algemeen geldt voorde productie van amorf silicium (a-Si), dat het verhogen van de waterstof-verdunning van het brongas silaan, de kwaliteit van het materiaal positiefbeïnvloedt, totdat het materiaal nanokristallijn wordt. Net voor deze over-gang vertoont het materiaal hoge interne spanning, wat tot gevolg kan hebbendat de siliciumlagen losspringen van het substraat. In het nc-Si gebied kan dekristalfractie van het materiaal beïnvloed worden door het veranderen van dewaterstofverdunning en door het veranderen van het toegevoerde vermogen.

Met deze siliciumlagen, samen met gedoteerde siliciumlagen gedeponeerdbij lage temperatuur, fabriceerden we a-Si zonnecellen met een intrinsieke laagvan 275 nm, zowel op glassubstraten als op PC substraten. Omdat over hetalgemeen siliciumlagen gedeponeerd bij lage temperatuur niet zo goed vankwaliteit zijn als lagen gedeponeerd bij optimale (hogere) temperaturen, kanrecombinatie van door licht gegeneerde ladingsdragers de werking van zonne-cellen verslechteren, door een lagere open-klemspanning en vulfactor. Dezeproblemen kunnen verminderd worden door een dunnere laag intrinsiek si-licium te gebruiken, maar hierdoor wordt een goede lichtopsluitingstechniekonmisbaar. Voor a-Si hebben we verschillende lichtopsluitingstechnieken gesi-muleerd en experimenteel getest, door gebruik te maken van geperste structu-ren in PC en willekeurige piramidestructuren op glas: Regelmatige piramides,veel groter dan de effectieve golflengte van zichtbaar licht op PC (micropirami-des), regelmatige piramidestructuren vergelijkbaar met de effectieve golflengtevan licht op PC (nanopiramides) en piramidestructuren op nanoschaal, zo-

Samenvatting 121

als gebruikelijk is voor commerciële zonnecelproductie (Asahi U-type). Beidestructuren op PC zorgen voor een verhoogde lichtabsorptie van de cellen. Metmicropiramide substraten behaalden we een initiële conversie-efficiëntie van6.4% en 7.4% op nanopiramide substraten, beide op PC. Cellen gedeponeerdonder dezelfde omstandigheden op Asahi U-type hadden een initiële efficiën-tie van 7.6%. Nanokristallijn silicium cellen met een ruwgeëtste aluminiumgedoteerde zinkoxidelaag (ZnO:Al) op glassubstraten vertoonden een initiëleefficiëntie van 6.9%, met een intrinsieke absorptielaag van slechts 750 nm.

Nadat we de dikte van de intrinsieke laag van de nc-Si cel hadden geop-timaliseerd, vertoonde een combinatie van een a-Si cel en een nc-Si zonnecelin een tandemstructuur (het zogenaamde micromorph concept), gedeponeerdbij 130°, een initiële efficiëntie van 9.5%, op ruwgeëtste ZnO:Al. Wanneerdezelfde fabricagemethode werd gebruikt om cellen aan te groeien op gestruc-tureerd PC, bleek dat de kristalgroei van de nc-Si cel zich anders gedroeg, watresulteerde in a-Si/a-Si tandemstructuren. Door de waterstofverdunning vande de nc-Si intrinsieke laag te veranderen konden we toch een micromorphe celmaken. De lichtopsluiting in de nc-Si deelcel van de tandemcel op gestructu-reerd PC werkte minder goed dan die van de tandemcel op ruwgeëtste ZnO:Al.

List of Publications

Publications within the scope of this thesisM.M. de Jong, J. de Koning, J.K. Rath and R.E.I. Schropp, Identificationof various plasma regimes in very high frequency PECVD of amorphous andnanocrystalline silicon near the phase transition, Proceedings of the 25thEUPV SEC Conference, V alencia, Spain, 3149-3151, 2010.

M.M. de Jong, A. Mohan, J.K. Rath, and R.E.I. Schropp, Temperature depen-dence of the ion energy distribution in a hydrogen diluted silane VHF plasma,AIP Conference Proceedings, 1397, 411-412, 2011.

M.M. de Jong, J.K. Rath, R.E.I. Schropp, P.J. Sonneveld, G.L.A.M. Swinkels,H.J. Holterman, J. Baggerman and C.J.M. van Rijn, Geometric light confine-ment in a-Si thin film solar cells on micro-structured substrates, Proceedingsof the 26th EUPV SEC Conference, Hamburg, Germany, 370-372, 2011.

M.M. de Jong, J. De Koning, J.K. Rath and R.E.I. Schropp. An opticalanalysis tool for avoiding dust formation in very-high frequency hydrogen di-luted silane plasmas at low substrate temperatures, Physics of P lasmas, 19,020703, 2012.

M.M. de Jong, J.K. Rath, R.E.I. Schropp, P.J. Sonneveld, G.L.A.M. Swinkels,H.J. Holterman, J. Baggerman, C.J.M. van Rijn and E.A.G. Hamers, A novelstructured plastic substrate for light confinement in thin film silicon solar cellsby a geometric optical effect, Journal of Non−Crystalline Solids, 358(17),2308-2312, 2012.

M.M. de Jong, J.K. Rath and R.E.I. Schropp, Very thin micromorph tan-dem solar cells deposited at low substrate temperature, Materials Research

124 List of Publications

Society Symposium Proceedings, 1426, 45-49, 2012.

M.M. de Jong, J. Baggerman, C.J.M. van Rijn, P.J. Sonneveld, G.L.A.M.Swinkels, H.J. Holterman, J.K.Rath and R.E.I. Schropp, Scattering, diffrac-tion and geometric light trapping in thin film amorphous silicon solar cellson plastic substrates, Materials Research Society Symposium Proceedings,1426, 155-160, 2012.

M.M. de Jong, P.J. Sonneveld, J. Baggerman, C.J.M. van Rijn, J.K. Rath andR.E.I. Schropp, Utilization of geometric light trapping in thin film silicon solarcells: Simulations and experiments, Progress in Photovoltaics, publishedonline, DOI: 10.1002/pip.2299

Publications outside the scope of this thesisJ.K. Rath, M.M. de Jong, A. Verkerk, M. Brinza and R.E.I. Schropp, Gasphase conditions for obtaining device quality amorphous silicon at low tem-perature and high deposition rate, Materials Research Society SymposiumProceedings, 1153, 463-468, 2009.

A.D. Verkerk, M M. de Jong, J.K. Rath, M. Brinza, R.E.I. Schropp, W.J.Goedheer, V.V. Krzhizhanovskaya, Y.E. Gorbachev, K.E. Orlov, E.M. Khilke-vitch and A.S. Smirnov, Compensation of decreased ion energy by increasedhydrogen dilution in plasma deposition of thin film silicon solar cells at lowsubstrate temperatures, Materials Science and Engineering B : Solid − StateMaterials for Advanced Technology, 159-160(C), 53-56, 2009.

J.K. Rath, Y. Liu, M.M. de Jong, J. De Wild, J.A. Schuttauf, M. Brinza andR.E.I. Schropp, Transparent conducting oxide layers for thin film silicon solarcells, Thin Solid F ilms, 518(24SUPPL.), e129-e135, 2010.

Nawoord

Nou, het is af. Juhluh!

Het schrijven van een proefschrift is veel werk. Gelukkig heb ik het nietalleen hoeven doen. Daarom wil ik graag een aantal mensen bedanken.

Allereerst dank ik mijn begeleiders. Jatin, dank voor de dagelijkse bege-leiding, voor je rijke kennis als het gaat om halfgeleiders en voor je positievekijk op de wereld, maar vooral op wetenschappelijke resultaten. Ik vond hetprettig dat je altijd beschikbaar was voor advies. Ruud ben ik vooral dank-baar dat hij me de gelegenheid gaf om aan dit project te beginnen en voorhet nauwkeurig lezen en verbeteren van alle teksten, plaatjes en praatjes dieik heb geproduceerd de afgelopen jaren.

Onderzoek met een sterk technische ondertoon kan niet zonder technici.Ik ben dan ook mijn dank verschuldigd aan Bart, Martin, Karine, Caspar,Gerard, Ruurd en Roberto voor het geduldig maken van talloze laagjes (zekerbij voorstellen als ’Misschien moeten we deze serie nog een keer doen’), oplossenvan vacuümproblemen, op peil houden van de characterisatietools en hun kijkop praktische problemen.

Piet (en later Jim en Theo), Gert-Jan en Henk Jan uit Wageningen dank ikvoor de samenwerking. De inkijkjes in de wondere wereld van de kastuinbouwvond ik intrigerend. Cees en Jacob dank ik voor de verschillende substraten.Zonder had ik een groot deel van mijn onderzoek niet uit kunnen voeren.Het werk van mijn studenten die ik begeleid heb: Jaap, Rob, Robin, Yaldaen in zekere zin kleine Casper hebben allemaal een plekje in het proefschriftgevonden, waarvoor dank.

Uiteindelijk ben ik lang aanwezig geweest in de groep ’Physics of Devices’.Dat is geen toeval. Ik heb het altijd een leuke groep mensen gevonden. Alsdunnefilmgroentje werd ik aan de arm genomen door Arjan, Monica en Hon-gbo. Met ’generatiegenoten’ Yanchao, Jan-Willem, Jessica, Sylvester, RuudBen Diederick heb ik veel plezier gehad, op het werk, maar ook op conferenties

126 Nawoord

en daarbuiten. Verder hoop ik dat Kuang, Henriette, Xin, Zachar, Kees, Pim,Akshatha, Caterina, Oumkelthoum, Lourens, Marcel en Wilfried net zoveelplezier zonder me hebben. De lol was er voor mij even af toen duidelijk werddat we als groep naar Eindhoven moesten verdwijnen, maar ik hoop dat julliedaar ook je draai vinden. Zoals Riny laatst zei: ’Het was gewoon een heel ge-zellig zootje en dat mis ik.’ Riny, dank voor je gezelligheid en levenswijsheid.Ik denk dat het voor mij ook zo gaat zijn.

Curriculum Vitae

The author was born on March 6, 1981 in Laren (NH), the Netherlands. Heobtained his secondary school degree in 1999 from ’Gymnasium Celeanum’ inZwolle. From 2000 to 2008 he studied Physics at Utrecht University, fromwhich he graduated in 2008 with a master degree titled ’Nanomaterials: Che-mistry and Physics’. For his master research he carried out research on the ionenergy distributions in silane and hydrogen deposition plasmas in the group’Physics of Devices’ at Utrecht University. In the same group, he started hisPhD research under the supervision of Prof. dr. R.E.I. Schropp and Dr. J.K.Rath on the topic of thin film silicon solar cells deposited at low depositiontemperatures and direct deposition of these cells on plastic substrates. Severallight trapping techniques were investigated. The results are presented in thisthesis.

light trapping in thin film silicon solar cells on plastic substrates

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