Advances in Photovoltaics: Part 3 SERIES EDITORS

VOLUME NINETY

SEMICONDUCTORS AND SEMIMETALS Advances in Photovoltaics: Part 3 SERIES EDITORS

EICKE R. WEBER Director Fraunhofer-Institut fur€ Solare Energiesysteme ISE Vorsitzender, Fraunhofer-Allianz Energie Heidenhofstr. 2, 79110 Freiburg, Germany CHENNUPATI JAGADISH Australian Laureate Fellow and Distinguished Professor Department of Electronic Materials Engineering Research School of Physics and Engineering Australian National University Canberra, ACT 0200 Australia VOLUME NINETY

SEMICONDUCTORS AND SEMIMETALS Advances in Photovoltaics: Part 3

Edited by

GERHARD P. WILLEKE Fraunhofer Institute for Solar Energy Systems ISE, Freiburg, Germany

EICKE R. WEBER Fraunhofer Institute for Solar Energy Systems ISE, Freiburg, Germany

AMSTERDAM • BOSTON • HEIDELBERG • LONDON NEW YORK • OXFORD • PARIS • SAN DIEGO SAN FRANCISCO • SINGAPORE • SYDNEY • TOKYO Academic Press is an imprint of Elsevier Academic Press is an imprint of Elsevier 32 Jamestown Road, London NW1 7BY, UK 525 B Street, Suite 1800, San Diego, CA 92101-4495, USA 225 Wyman Street, Waltham, MA 02451, USA The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, UK

First edition 2014

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ISBN: 978-0-12-388417-6 ISSN: 0080-8784

For information on all Academic Press publications visit our website at store.elsevier.com CONTENTS

Contributors vii

1. State-of-the-Art Industrial Crystalline Silicon Solar Cells 1 Giso Hahn and Sebastian Joos

1. Introduction 4 2. Operation Principle of a c-Si Solar Cell 10

3. The Basic Firing Through SiNx:H Process 19 4. Recent Developments on Solar Cell Front Side 34 5. Advanced Emitter Formation 40 6. Industrial PERC-Type Solar Cells 51 7. Summary and Outlook 60 Acknowledgments 62 References 62 2. Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 73 Christophe Ballif, Stefaan De Wolf, Antoine Descoeudres, and Zachary C. Holman

1. Introduction 74 2. Passivating c-Si Surfaces with a-Si:H 76 3. From Passivated Wafers to Complete Solar Cells 83 4. Losses in Silicon Heterojunction Solar Cells 95 5. Industrialization and Commercialization 99 6. Future Directions and Outlook 108 Acknowledgments 110 References 110 3. Overview of Thin-Film Solar Cell Technologies 121 Bernhard Dimmler

1. Introduction 121 2. Market Shares of TF in PV 123 3. TF Device Efficiencies in Laboratory and Industry 125 4. Future Developments of TF Technologies in PV 128 References 136

Index 137 Contents of Volumes in this Series 141

v This page intentionally left blank CONTRIBUTORS

Christophe Ballif Photovoltaics and Thin-Film Electronics Laboratory, Institute of Microengineering (IMT), Ecole Polytechnique Fe´de´rale de Lausanne (EPFL), Neuchâtel, Switzerland. (ch2) Stefaan De Wolf Photovoltaics and Thin-Film Electronics Laboratory, Institute of Microengineering (IMT), Ecole Polytechnique Fe´de´rale de Lausanne (EPFL), Neuchâtel, Switzerland. (ch2) Antoine Descoeudres Photovoltaics and Thin-Film Electronics Laboratory, Institute of Microengineering (IMT), Ecole Polytechnique Fe´de´rale de Lausanne (EPFL), Neuchâtel, Switzerland. (ch2) Bernhard Dimmler Manz AG, Reutlingen, Germany. (ch3) Giso Hahn Department of Physics, University of Konstanz, Konstanz, Germany. (ch1) Zachary C. Holman School of Electrical, Computer, and Energy Engineering, Arizona State University, Tempe, Arizona, USA. (ch2) Sebastian Joos Department of Physics, University of Konstanz, Konstanz, Germany. (ch1)

vii This page intentionally left blank PREFACE

The rapid transformation of our energy supply system to the efficient use of renewable energies remains to be one of the biggest challenges of mankind that increasingly offers exciting business opportunities as well. This truly global-scale project is well on its way. Harvesting solar energy by photovoltaics (PV) is considered to be a cornerstone technology for this transformation process. This book presents the third volume in the series “Advances in Photovoltaics” in Semiconductors and Semimetals. This series has been designed to provide a thorough overview of the underlying physics, the important materials aspects, the prevailing and future solar cell design issues, production technologies, as well as energy system integration and characterization issues. In this volume, three distinctly different solar cell technologies are covered in detail, ranging from state-of-the-art crystalline silicon technology, the workhorse of the booming PV market, to one of the most advanced technologies, silicon heterojunction cells, and to an overview of thin film solar cell technologies. Therefore, this volume represents a cornerstone of “Advances in Photovoltaics,” as the first and the third chapter together cover more than 98% of the current PV world market volume. The second chapter provides a glimpse into the future of highly efficient crystalline Si PV technologies that will allow further decrease in the cost of PV-generated electricity available from premium modules with top performance produced at prices that will become competitive with present-day low-cost PV modules. Following the tradition of this series, all chapters are written by world-leading experts in their respective field. In the past 2 years, since the introduction to the first volume of this series has been written, the world PV market has undergone a decisive transformation. Huge production overcapacity, established especially in Asia, resulted in rapidly declining prices, often to values beyond the production costs, when fire sales of module supplies were the only way to generate desperately needed cash for financially stressed companies. Subsequently, many companies went into insolvency, followed by either restructuring under new ownership, often from abroad, or a complete shutdown of the production lines. The PV equipment manufacturers were especially hard hit, as they had to survive several years practically without any new orders.

ix x Preface

Today we experience a new development: decreasing global production capacity begins to meet further increasing PV market size, the growth of which is fueled worldwide by the low cost of solar electricity. The consequence of this process will be the further decentralization of electricity supply, as PV systems increasingly allow owners of homes and industry to produce electricity on their own roofs and free areas, to the benefit of energy independence and the world climate, that desperately needs rapid further market penetration of renewables to decrease the emission of climate gases.

GERHARD P. WILLEKE AND EICKE R. WEBER Fraunhofer ISE, Freiburg, Germany CHAPTER ONE

State-of-the-Art Industrial Crystalline Silicon Solar Cells

Giso Hahn1, Sebastian Joos Department of Physics, University of Konstanz, Konstanz, Germany 1Corresponding author: e-mail address: [email protected]

Contents

1. Introduction 4 1.1 History 4 1.2 General routes for cost reduction 5 1.3 PV market today 7 1.4 Basic structure of an industrial c-Si solar cell 9 2. Operation Principle of a c-Si Solar Cell 10 2.1 Band diagram 10 2.2 Solar cell parameters 12 2.3 Fundamental efficiency limit of an ideal c-Si solar cell 13 2.4 Two-diode model 14 2.5 Radiative recombination 14 2.6 Auger recombination 15 2.7 SRH recombination 16 2.8 Surface recombination 17 2.9 Recombination and saturation current density 18 2.10 Optical losses 18 3. The Basic Firing Through SiNx:H Process 19 3.1 Wafer washing, texturization, and cleaning 20 3.2 Phosphorus diffusion 22 3.3 Edge isolation 25 3.4 SiNx:H deposition 25 3.5 Metallization via screen-printing 27 3.6 Solar cell characterization 33 4. Recent Developments on Solar Cell Front Side 34 4.1 Wafer sawing 34 4.2 Alkaline wafer texturing 35 4.3 Front contact metallization 35 5. Advanced Emitter Formation 40 5.1 Improvement of homogeneous emitters 41 5.2 Selective emitters 42 6. Industrial PERC-Type Solar Cells 51 6.1 Dielectric rear side passivation 52

6.2 Formation of local rear contacts 54 6.3 Boron–oxygen related degradation 57 6.4 State-of-the-art industrial PERC solar cells 59 7. Summary and Outlook 60 Acknowledgments 62 References 62

ABBREVIATIONS A area ALD atomic layer deposition APCVD atmospheric pressure chemical vapor deposition ARC antireflective coating a-Si amorphous silicon BSF back surface field Bs substitutional boron concentration cA,n (cA,p) Auger recombination coefficient for electrons (holes) crad radiative recombination coefficient c-Si crystalline silicon Cz Czochralski d layer/wafer thickness dBSF + D diffusion constant in the BSF DI deionized Dn (Dp) diffusion constant of electrons (holes) E energy ECV electrochemical capacitance voltage EF (EFi) (intrinsic) Fermi energy level EFG edge-defined film-fed growth EFn (EFp) quasi-Fermi energy level of electrons (holes) Eg band gap energy Ephot photon energy EQE external quantum efficiency Et energetic position of the trap level EVA ethylene vinyl acetate FCA free carrier absorption FF fill factor FZ float zone h Planck’s constant HIT heterojunction with intrinsic thin-layer I current IBC interdigitated back contact IPA isopropyl alcohol IQE internal quantum efficiency j current density j0 saturation current density j01 ( j02) saturation current density of the first (second) diode j0e saturation current density of the emitter State-of-the-Art Industrial Crystalline Silicon Solar Cells 3

jl light-generated current density jsc short circuit current density k Boltzmann’s constant + L diffusion length in the BSF LFC laser fired contacts Ln (Lp) diffusion length of electrons (holes) LPCVD low pressure chemical vapor deposition mono-Si monocrystalline Si mpp maximum power point mc-Si multicrystalline Si n electron concentration + ++ n (n ) (very) highly n-doped n0 electron concentration in the dark NA (ND) acceptor (donor) concentration + NA acceptor concentration in the BSF nair (nSi, nSiN) refractive index of air (c-Si, SiN) ni intrinsic carrier concentration Nt trap density Nts areal trap density at the surface Oi interstitial oxygen p hole concentration + p highly p-doped p0 hole concentration in the dark PECVD plasma-enhanced chemical vapor deposition PERC passivated emitter and rear cell PERL passivated emitter and rear locally diffused PERT passivated emitter and rear totally diffused pphot photon power density PSG phosphor silicate glass Psurf phosphorous surface concentration Ptot total power loss PV photovoltaic q elementary charge R recombination rate RA Auger recombination rate Rrad radiative recombination rate Rs series resistance Rs,tot total series resistance RSRH Shockley-Read-Hall recombination rate Rsh shunt resistance Rsheet sheet resistance of the emitter s (sn)(sp) surface recombination velocity (of electrons or holes) sb surface recombination velocity at the backside SCR space charge region seff effective surface recombination SIMS secondary ion mass spectrometry SRH Shockley-Read-Hall STC standard test conditions (1000 W/m2, AM1.5g spectrum, 25 C) UMG upgraded metallurgical grade 4 Giso Hahn and Sebastian Joos

V voltage vn (vp) thermal velocity of electrons (holes) Voc open circuit voltage Wp Watt peak (power of 1 W under STC) α absorption coefficient ΔEF splitting of quasi-Fermi levels Δn excess charge carrier density η conversion efficiency Φ photon flux λ wavelength ρSi density of Si ρ resistivity σn (σp) capture cross section for electrons (holes) τA Auger lifetime τb bulk lifetime τeff effective lifetime τrad radiative lifetime τSRH Shockley, Read, Hall lifetime τ minority charge carrier lifetime

1. INTRODUCTION

Solar cells fabricated based on crystalline Si (c-Si) generate electricity from sunlight by absorbing photons and generating electron–hole pairs, which are separated by a pn-junction. The pn-junction creates an electric field in the semiconductor and the separated charge carriers have to leave the solar cell via electrical contacts to perform work in an external circuit. A solar cell in operation is therefore essentially an illuminated large area diode, where emitter and base regions are contacted by metals to extract the carriers.

1.1. History The first c-Si solar cell operating using the principle described above was reported in 1953 (Chapin et al., 1954), although research toward this achievement dates back to the 1940s (e.g., Ohl, 1941; Shockley, 1950). In the decades to follow, research was first directed toward application of the photovoltaic (PV) effect in space (powering satellites) or for terrestrial stand-alone systems. As for those applications the total cost of power generation was not the main issue, research was mainly driven by improving the conversion efficiency η, which is the ratio between output power from the PV device (generated from the solar cell or complete solar module) and State-of-the-Art Industrial Crystalline Silicon Solar Cells 5 input power (impinging photon flux). The oil crisis in 1973 led to considerations to use PV also for terrestrial applications in larger scale as an alternative to fossil fuels. Since then a lot of R&D activities was focused on reducing the cost of PV electricity generation to make it attractive for market penetration. In research, a lot of progress was made in improving efficiency by developing new cell designs and applying novel processing steps, leading to efficiencies as high as 25% using standard test conditions (STC: 1000 W/m2 illumination, AM1.5g spectrum, 25 C) in 1999 (Zhao et al., 1999), indicating the efficiency potential of c-Si. This efficiency was reached on extremely pure float zone (FZ) silicon and on small scale (4 cm2) without the main part of the front side metallization grid being taken into account for the efficiency measurement (so-called designated area measurement) and using a very complex processing scheme. For most industrial applications, a full area measurement and cost-effective c-Si materials are of higher interest. In addition, the number and complexity of processing steps needed for cell fabrication has to be low, to allow a cost-efficient production. Here, the main challenge for industrial c-Si solar cells becomes visible: there is a trade-off between more complex processing on higher quality material allowing higher efficiencies, and less complex processing, e.g., in combination with a lower c-Si material quality.

1.2. General routes for cost reduction The lower efficiency for lower cost materials and less complex processing might be advantageous cost-wise at cell level, but as there are also area related cost factors at module and system level (e.g., costs for module glass and installation), the question which route is more promising is not easy to answer. Therefore, a lot of different technologies have been developed over the past decades. This includes c-Si materials as well as solar cell fabrication processes. The Si feedstock of highest quality stems from the so-called Siemens route using rods for Si production from the gas phase, which still accounts for the majority of produced Si wafers for industrial solar cells, with fluidized bed reactors as an alternative (Fabry and Hesse, 2012). So-called upgraded metallurgical grade (UMG) Si can be produced with significantly less energy needed per kg of fabricated Si, but a higher impurity concentration is the consequence, with relatively high amounts of, amongst others, B and P still present acting as doping elements in Si. This might cause problems as after crystallization the material will be partly compensated, and due to 6 Giso Hahn and Sebastian Joos different segregation coefficients of B and P their concentrations and therefore resistivity, influenced by the net doping, changes with ingot height (Ceccaroli and Pizzini, 2012; Heuer, 2013). For c-Si materials, three different material classes have been important for PV in the past, as they have already been in industrial production in significant quantities. Monocrystalline Si (mono-Si) pulled using the Czochralski (Cz) method shows the lowest amount of extended crystal defects (like, e.g., grain boundaries, dislocations, precipitates), but normally contains a high amount of O, mainly in interstitial form (Oi)(Zulehner, 1983). Cast multicrystalline Si (mc-Si) can be produced in a more cost-effective way, but contains due to the crystallization method used a higher amount of extended crystal defects and impurities in interstitial or precipitated form, originating mainly from the crucible wall and the crucible coating (Buonassisi et al., 2006; Schubert et al., 2013). See Coletti et al. (2012) for an overview on the role of impurities in c-Si for solar cells. For both methods, the crystallized ingot has to be sliced in wafers for subsequent solar cell processing. To avoid kerf and other Si material losses that easily amount to >50%, ribbon-Si techniques have been developed, crystallizing the Si wafer directly from the Si melt (Hahn and Schonecker,€ 2004). Of the three technology groups, ribbon Si is the most cost-effective technique to produce wafers, but these wafers normally show the highest defect densities, reducing the electronic quality of the as-grown wafer. Apart from Si wafer quality, solar cell process complexity is the other main parameter determining the efficiency and cost structure of the solar cell. In this contribution, focus is laid on industrial solar cell production, but for a more complete picture also PV module and system aspects should be considered. The heart of a solar module and every PV system is the solar cell. The cells are stringed in series so that the same amount of current flows through all cells in a string and the voltages of the cells add up. This makes proper sorting of cells a necessity to ensure that cells of similar performance end up in a string, as the cell with the lowest current at operation conditions determines the current flowing through the string. Therefore, for all cells not only the peak efficiency, but also a tight distribution of cell parameters is important to facilitate sorting and matching of the cells. This means that in industrial fabrication homogeneous Si wafer quality and stable processes with large process windows are desired to minimize the spread of quality in c-Si solar cell production. In this chapter, an overview on industrial state-of-the-art c-Si solar cells is given. As there is not only one industrial solar cell process, but a variety of different processes applied for different cell designs, we will restrict the State-of-the-Art Industrial Crystalline Silicon Solar Cells 7 overview on the most common cell architectures. Other cell designs already used in industrial scale such as the interdigitated back contact (IBC), commercialized by company SunPower Corp. (Cousins et al., 2010), or the heterojunction with intrinsic thin-layer (HIT) concept pioneered by Sanyo (now Panasonic) (Ballif et al., 2014) allow for the highest efficiencies in commercial c-Si solar cells on large area cells with lab cell record efficiencies up to 25% on large area cells (Smith et al., 2014; Taguchi et al., 2013) and even 25.6% with a combined IBC-HIT approach (Panasonic, 2014), but the processes differ significantly from mainstream technology. Therefore, these designs of very highly efficient c-Si solar cells will be treated in other chapters (e.g., Ballif et al., 2014).

1.3. PV market today Figure 1.1 demonstrates the very dynamic growth of commercial PV over the past decades, spanning more than four decades from around 1 MWp1 in the early 1970s to >30 GWp in 2011. Annual growth rates over the past 10 years have been in the order of 50%, mainly driven by market stimulation programs like, e.g., the renewable energy law with a guaranteed feed-in tar- iff in Germany. As the German feed-in tariffs have been adjusted recently and the German PV market was the strongest worldwide, the growth slowed down in 2012 and 2013. Strong growth in recent years allowed for a tremendous reduction in production cost due to scaling effects in mass production

10,000

1000 er (MWp)

100

10 PV-module pow PV-module

1975 1980 1985 1990 1995 2000 2005 2010 Figure 1.1 Yearly production/shipment of solar modules. Data from PV News, Photon, and Mehta (2014).

1 Watt peak (Wp) refers to the power generated under STC. 8 Giso Hahn and Sebastian Joos as well as new and optimized processing technologies. This so-called learning curve effect of PV resulted in an average module price reduction of around 20% for every doubling of cumulated PV production (Nemet and Husmann, 2012). The continuing reduction in processing costs results in costs of a kWh generated by PV being now in the range of electricity generated from fossil fuels (depending on the installation site) (Kost et al., 2013). The market share of different PV technologies shown in Fig. 1.2 reveals that c-Si still shows by far the highest market penetration, with thin film technologies like amorphous Si (a-Si), CdTe and CuInxGa(1x)Se2 (CIGS) not really gaining market share above a 10–15% level. In contrast, latest figures indicate an even further increasing market share for c-Si of 90% in 2013, with roughly 67% based on mc-Si and 23% on mono-Si (Mehta, 2014). It is interesting to note that mono-Si lost market share to mc-Si in the past decade. This can be explained by the huge production expansion programs happening at most PV manufacturers in the past, as mc-Si technology seems to be easier to ramp up and was the more cost- effective way of production in the past. Whether this will hold true in the future, with new cell designs allowing for higher efficiency approaching the market, remains to be seen. The market share of ribbon-Si dropped to almost zero as the two main technologies edge-defined film-fed growth (EFG) and string ribbon are no longer on the market, due to the disappearing of their production companies Schott Solar and Evergreen Solar as well as EverQ, respectively.

100 90 80 Others CIGS 70 CdTe a-Si 60 Ribbon-Si 50 Multi-Si Mono-Si 40

Technology(%) 30 20 10 0 1997 1998 1999 2000 2001 2002 2003 2004 2005 2006 2007 2008 2009 2010 2011 Figure 1.2 Market share of different PV technologies. Data from PV News and Photon. State-of-the-Art Industrial Crystalline Silicon Solar Cells 9

1.4. Basic structure of an industrial c-Si solar cell A schematic of the basic structure for a typical state-of-the-art industrial c-Si solar cell is shown in Fig. 1.3. The base is p-type material, moderately 16 3 B doped to a resistivity of around 1 Ω cm (NA ¼1.510 per cm ). The ++ 2 20 emitter is n -doped using P with high surface concentration ND >10 per cm3, and the front surface is textured to allow a better incoupling of impinging photons (lower reflectivity). The emitter is covered by a thin dielectric layer of H-rich silicon nitride (SiNx:H), acting as antireflective coating (ARC), surface passivation layer, and reservoir of H. On the front, the metallization finger grid is realized by Ag paste, fired through the SiNx:H layer at high temperature. On the rear, a full area contact is realized by Al paste, which forms an alloy with Si during the firing step, resulting in an Al doped p+-region (around 1019 per cm3) at the rear after cool down to room temperature (back surface field, BSF). To allow interconnection of the individual cells for module integration using soldering, stripes or pads of Ag/Al paste are used at the rear side, as Al is not solderable. The complete cell thickness is around 180 μm (note that features shown in Fig. 1.3 are not to scale). The formation of the respective regions of the cell will be dealt with in more detail in the following sections. The use of H-rich SiNx:H layers for PV (Morita et al., 1982)inthe so-called “firing through SiNx:H process” has been pioneered by Kyocera (Kimura, 1984; Takayama et al., 1990) and Mobile Solar for their EFG ribbon-Si material (Cube and Hanoka, 2005). In the 1990s, other companies and research institutes like, e.g., IMEC (Szlufcik et al., 1994) and others developed the process further. The breakdown of costs for c-Si module production in Fig. 1.4 reveals that wafer and module costs are the dominating factors.

hν

SiNx:H Ag

n+ p-Si Electron Hole p+

Al Figure 1.3 Schematic basic structure of an industrial c-Si solar cell in cross section (not to scale).

2 The superscripts + and ++ indicate a high and a very high doping concentration, respectively. 10 Giso Hahn and Sebastian Joos

Wafer Cell production Module

26% 36%

38% Figure 1.4 Breakdown of c-Si PV module manufacturing costs. Data from Goodrich et al. (2013).

Excellent early (e.g., Szlufcik et al., 1997) and more recent (e.g., Gabor, 2012; Neuhaus and Mu¨nzer, 2007) review papers on low-cost industrial c-Si solar cell fabrication exist, forming the base of this chapter. Since then new technologies have emerged, allowing for a reduction of costs as well as efficiency losses and therefore an increase of efficiency in mass production. To tackle these losses, the next section will describe the physics involved in the operation principle of a solar cell.

2. OPERATION PRINCIPLE OF A c-SI SOLAR CELL 2.1. Band diagram The fundamental operation principle of a c-Si solar cell is visualized in the band diagram shown in Fig. 1.5. The doping gradient due to the abrupt change in doping concentration at the pn-junction results in electrons (free majority carriers in the n-region) diffusing from the n-region into the p-region and holes (free majority carriers in the p-region) diffusing into the n-region. The remaining ionized doping atoms at lattice sites (positively charged in the n-region, negatively charged in the p-region) form the space charge region (SCR) extending into both sides of the pn-junction. The electric field hinders the free carriers to completely diffuse into the regions of opposite doping, when equilibrium between diffusion and drift current of free carriers is reached. The built-up electric field causes bending of the energy bands, with the Fermi energy EF as defined by the Fermi–Dirac function at a constant level (a horizontal line) in both regions. Upon illumination, absorbed photons excite electrons from the valence band to the conduction band via the internal photoelectric effect. State-of-the-Art Industrial Crystalline Silicon Solar Cells 11

Energy E Electron Conduction band

hν E F

E F

Valence band Hole

Metal p-type Si SCR n-type Si Metal Figure 1.5 Schematic band diagram of a c-Si solar cell with pn-junction, space charge region (SCR), photon absorption, charge carrier generation, and separation. Quasi-Fermi levels and EF in the metal contacts are indicated as well.

Absorption of one photon therefore generates an electron–hole pair, as the missing electron in the valence band is referred to as a hole. Free electrons and holes can diffuse until they recombine or reach the SCR. Here, charge carriers of different types are separated, electrons are accelerated into the n-region, holes into the p-region. In case of illumination, the semiconductor is not in thermal equilibrium anymore, and the relation for electron and hole concentrations n0 and p0, respectively, as defined for thermal equilibrium (without illumination or applied voltage)

n p ¼ n2 0 0 i , (1.1)

(with intrinsic carrier concentration ni) is not valid anymore and becomes E E np ¼ n2 Fn Fp > n2 i exp kT i , (1.2) with n and p being electron and hole concentrations, respectively. As both electron and hole concentrations are increased when the semiconductor is illuminated, two separate Fermi–Dirac functions for each carrier type have to be defined, with two resulting Fermi levels EFn and EFp referred to as quasi-Fermi levels of electrons and holes. Metal contacts with EF at roughly the same energetic position as for the majority carriers in the contacted Si region can extract carriers from both regions. The contact for the p-type region as depicted in Fig. 1.5 is ohmic, whereas the n-type contact is of Schottky-type (energy barrier for electrons). 12 Giso Hahn and Sebastian Joos

The barrier can be overcome via tunneling, provided it is thin enough and not too high.

2.2. Solar cell parameters An ideal solar cell can be described by a 1-diode model and the j–V characteristic of an illuminated diode qV j ¼ j j 0 exp kT 1 l, (1.3) with current density j, saturation current density j0, elementary charge q, Boltzmann’s constant k, and light-generated current density jl. j0 is defined as qD n2 qD n2 n i p i j0 ¼ + , (1.4) LnNA LpND with Dn (Dp) the diffusion constant of electrons (holes), NA (ND) the doping density of acceptors (donors) and Ln (Lp) the minority charge carrier diffusion length of electrons (holes). The resulting j–V curve is shown in Fig. 1.6. The maximum current density at V¼0 is the short circuit current density jjscj¼jl. The point of maximum power density (mpp) is also indicated, with the fill factor FF defined as

Current density/power density

Dark curve

Output power Illuminated curve V V mpp oc Voltage Open-circuit voltage

Maximum power point (MPP) j mpp j sc Short circuit current density Figure 1.6 Dark and illuminated j–V curve of a solar cell as well as output power in dependence of voltage. State-of-the-Art Industrial Crystalline Silicon Solar Cells 13

j V FF ¼ mpp mpp , (1.5) jscVoc resulting with the impinging photon power density pphot of photons with energy Ephot in the efficiency j V FF η ¼ sc oc : (1.6) pphot

2.3. Fundamental efficiency limit of an ideal c-Si solar cell In a semiconductor with band gap Eg (1.12 eV at 25 C for c-Si), photons with energy E>Eg can be absorbed, creating electron–hole pairs, while photons with E

ν ν ΔE h h F qVmpp Energy

1. 2. 3. 4. Figure 1.7 Fundamental loss mechanisms for an ideal pn-junction based solar cell. 1.

Transmission Ephot Eg, 3. Quasi-Fermi level splitting ΔEF

ΔEF

2.4. Two-diode model A real solar cell can be described by an equivalent circuit containing two diodes, with the addition of series resistance Rs, shunt resistance Rsh and a second diode accounting for recombination in the SCR with an ideality factor generally assumed to be 2 (Fig. 1.8). qVðÞ jRs qVðÞjRs ðÞV jRs j ¼ j01 exp 1 + j02 exp 1 + jl: kT 2kT Rsh (1.7)

Contributions to Rs are ohmic resistive losses in emitter, base, and metallization as well as the contact resistance between semiconductor and metal. Finite Rsh values are caused by alternative current paths short circuiting the diode (e.g., around the cell’s edge, by a damaged emitter or current paths through the SCR). Apart from ohmic losses, recombination of generated charge carriers can occur, limiting performance of the solar cell.

2.5. Radiative recombination Radiative recombination refers to direct band-to-band transitions of an electron from the conduction band to the valence band while emitting a photon. It is the inverse process of photon absorption. The generated excess charge carrier density Δn with

n ¼ n0 + Δn and p ¼ p0 + Δn (1.8)

R j j S 01 02 j I

R Sh

Figure 1.8 Equivalent circuit of a real pn-junction solar cell. State-of-the-Art Industrial Crystalline Silicon Solar Cells 15 can be reduced due to recombination of charge carriers with a recombination rate R defining the lifetime τ of excess charge carriers Δn τ ¼ : R (1.9) c-Si is an indirect band gap semiconductor. In addition to an electron (in the conduction band) and a hole (in the valence band), a phonon is necessary for the band-to-band transition to occur due to conservation of momentum. Therefore, this mechanism is not probable and can normally be neglected in c-Si. With the radiative recombination coefficient crad, the net rate Rrad for this type of recombination becomes3 R ¼ c np n2 rad rad i , (1.10) resulting for low injection (Δn much lower than doping concentration4)in the radiative lifetime 1 τrad¼ (1.11) cradp0 for p-doped material.

2.6. Auger recombination Instead of creating a photon, the energy of the recombination process can be used to excite another existing free charge carrier (an electron in the conduction band or a hole in the valence band). This charge carrier thermalizes after excitation toward the band edge, converting the recombination energy into phonons. With the Auger recombination coefficients cA,n and cA,p for electrons and holes, respectively, the Auger recombination rate reads R ¼ c nnp n2 c pnp n2 : A A,n i + A,p i (1.12) As above, for low injection we obtain the Auger lifetime for p-doped material 1 τ ¼ : (1.13) A c p2 A,p 0

3 Note that we are only interested in the recombination rate of the excess charge carriers (therefore 2 npni , subtracting recombination occurring also in thermal equilibrium). 4 At room temperature, all dopants are assumed to be ionized (NA ¼p0 in p-type material), and therefore Δnp0 for low injection. 16 Giso Hahn and Sebastian Joos

Auger recombination as a three-particle process is only relevant for high doping concentrations >1017 per cm3 in standard industrial solar cells.

2.7. SRH recombination Energy levels in the band gap can trap free charge carriers and cause a very effective recombination mechanism, especially when their energetic position is close to mid-gap. This type of recombination was formulated by Shockley, Read, and Hall (Hall, 1952; Shockley and Read, 1952), using statistics of capture and emission of free carriers and is therefore referred to as SRH recombination. Its recombination rate 2 np ni RSRH ¼ (1.14) τpðÞn0 + n1 + Δn + τnðÞp0 + p1 + Δn with 1 1 Et EFi EFi Et τp ¼ , τn ¼ ,n1 ¼ niexp , p1 ¼ niexp , Ntvpσp Ntvnσn kT kT (1.15) includes the trap density Nt of the energy levels in the band gap, the thermal velocity of electrons and holes (vn, vp) and the capture cross sections of the trap for electrons and holes (σn, σp). Et is the energetic position of the trap level and EFi the position of the Fermi level in intrinsic c-Si. The SRH lifetime

τpðÞn0 + n1 + Δn + τnðÞp0 + p1 + Δn τSRH ¼ (1.16) p0 + n0 + Δn for p-type material (p0 n0), low injection (p0 Δn), and trap energy level at mid-gap (Et ¼EFi) reads 1 τSRH ¼ τn ¼ (1.17) Ntvnσn and is inversely proportional to the trap density as well as the thermal velocity and capture cross section of the minority carriers (electrons in p-type material). All recombination channels are acting in parallel, and the resulting bulk lifetime τb is given by State-of-the-Art Industrial Crystalline Silicon Solar Cells 17

1 1 1 1 ¼ + + : (1.18) τb τrad τA τSRH

2.8. Surface recombination At the crystal surface, dangling bonds5 are responsible for a multitude of defect levels distributed throughout the band gap. In analogy to the SRH recombination formalism in the bulk of the crystal, a lifetime of the charge carriers at the physical surface can be derived using areal instead of volume densities of charge carriers and traps. For p-type material in low injection, this results in

sn ¼ Ntsvnσn, (1.19) with the areal density of traps at the surface Nts, and sn being referred to as the surface recombination velocity s of electrons (minority carriers in p-type material) in units of cm/s. The influence of surface recombination on the observable effective lifetime can be expressed by a surface lifetime τs (Aberle, 1999)

1 1 1 1 2 ¼ + ¼ + α Dn, (1.20) τeff τb τs τb with α a solution of the transcendental equation (wafer thickness d) αd s tan ¼ , (1.21) 2 αDn which can be approximated with (Sinton and Cuevas, 1996)

d d2 τ : s + 2 (1.22) 2s Dnπ For reasonably good surface passivation with s <1000 cm/s, the second term can be neglected and 1 1 2s ¼ + : (1.23) τeff τb d

5 Dangling bonds are generally reconstructed bonds where the lengths and angles differ from their standard values in the c-Si bulk. 18 Giso Hahn and Sebastian Joos

2.9. Recombination and saturation current density

Recombination reduces the maximum current density jsc of the solar cell, as only minority charge carriers generated within roughly one diffusion length on either side of the pn-junction reach the junction and are injected into the region on the opposite side of the junction. But from Eq. (1.3) also strong influence of j0 on Voc can be seen, as for j¼0 kT jl kT jl Voc ¼ ln +1 ln : (1.24) q j0 q j0

As the diffusion lengths of both types of carriers in Eq. (1.4) are linked to recombination via the lifetime of the minority charge carriers pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi Ln,p ¼ Dn,pτeff , (1.25) maximizing the effective lifetimes in emitter and base is crucial for improving solar cell performance. Effective lifetime is affected by bulk lifetime and surface recombination velocity (Eq. 1.23), therefore good solar cells should combine a high τb (low recombination in bulk and emitter) and good surface passivation on emitter and base to reduce s.

2.10. Optical losses

If all impinging photons with Ephot >Eg were absorbed in the solar cell, with all of these photons contributing to the extracted current density, the max- 2 imum jsc would be around 44 mA/cm under STC. Apart from recombination losses described above, another fraction is lost due to optical losses. These losses include reflection at the front side (metal grid and ARC), absorption in the metal and ARC, absorption via free carrier absorption (FCA)6 and photons not being absorbed in c-Si (mostly long wavelengths photons7) leaving the cell. The different loss mechanisms are visualized in Fig. 1.9, where they are separated into optical and electrical losses.

6 Free carrier absorption is the absorption of a photon by an electron in the conduction band or a hole in the valence band without generation of additional free carriers. It is important in highly doped areas (emitter and BSF). 7 The absorption coefficient in c-Si with indirect bandgap leads to an absorption coefficient strongly varying with wavelength, leading for photons with wavelengths >1000 nm to absorptions lengths >200 μm. State-of-the-Art Industrial Crystalline Silicon Solar Cells 19

Shadowing loss (total reflection on metal) Incident photon flux F ARC reflection loss solar spectrum (mainly short wavelengths) AM1.5g Back reflection (mainly long wavelengths)

SiNx:H n+ Carrier loss emitter & SCR

ARC absorption loss Final carrier (mainly short wavelengths) Free carrier j q absorption flow sc/ Free carrier Carrier loss bulk p-Si absorption Carrier loss BSF BSF Al Rear absorption loss

Figure 1.9 Visualization of the conversion of photon flux into carrier flow in a standard industrial p-type Si solar cell with the optical and electrical losses as indicated.

3. THE BASIC FIRING THROUGH SiNx:H PROCESS As already mentioned in the introduction, most industrial solar cells today are fabricated based on a so-called “firing through SiNx:H” process (Fig. 1.3). Therefore, in this section we will describe this process in its basic form as it was developed in more detail (compare with, e.g., Neuhaus and Mu¨nzer, 2007; Szlufcik et al., 1997), before alternatives and improvements will be dealt with in the next sections. Generally, for every process step there are two options, inline or batch processing. Inline processing offers the possibility to fabricate solar cells with a minimum of handling steps and a smaller footprint due to the lack of storage room necessary for partially processed cells. On the other hand, not all processing steps can easily be performed inline and batch processing allows for more freedom in optimization. The first example of a complete true inline processing fabrication of solar cells was RWE Schott Solar’s SmartSolarFab in 2002. Nowadays, cell processing is normally done by a mixture of inline and batch processing equipment, as the throughput of machines used for the different steps is not the same. In addition, if single machines are not operational or have to be maintained, not the complete production is halted, but other parts within cell fabrication can continue to produce. Therefore, often several machines of the same type work in parallel to increase throughput and minimize the risk of bottlenecks. 20 Giso Hahn and Sebastian Joos

3.1. Wafer washing, texturization, and cleaning After crystallization, mono-Si and mc-Si wafers are sliced out of the Si ingot using wire saws, containing slurry with abrasives for cutting into the Si (Dold, 2014). This leaves, apart from contaminants, saw damage on both sides of the Si wafer with a depth in the range of up to 10 μm (depending on sawing conditions). After wafer washing, this saw damage has to be removed, as the disturbed region of the crystal (cracks, dislocations) is of poor electronic quality. For mono-Si, this is done in an alkaline wet chemical solution of KOH and isopropyl alcohol (IPA) at temperatures of around 80 C. The KOH solution etches the Si while the alcohol masks the surface randomly. Etching is anisotropic, with the result that the most densely packed crystal planes in c-Si have the slowest etch rate (the (111)-planes). If the wafer is (100)- oriented, the four (111) orientations in the diamond lattice of c-Si will randomly form square-based upright pyramids (Fig. 1.10). These pyramids very effectively reduce the reflectivity of the surface and therefore increase the incoupling of photons into c-Si. The etching reaction can be summarized as

3 Si + 2H2O+HO ! HSiO +2H2 (1.26) and consists of oxidation of Si, formation of a solvable salt and dissolving the salt in water (Neuhaus and Mu¨nzer, 2007). The surface is increased after random pyramid texturing by a factor of 1.7, which has consequences for surface passivation and saturation current densities of the emitter and the SCR. mc-Si does not offer a well-defined grain orientation at the wafer surface, as the grains are randomly distributed. Therefore, other texturing solutions had to be developed. Standard is an acidic solution based on HF and HNO3 without further additives (Einhaus et al., 1997; Hauser et al., 2003). The texture attacks the Si surface first at areas where not all Si bonds are perfectly saturated. Therefore, the saw damage is needed for a non-uniform attack of the surface. Existing surface defects like cracks are widened and a “worm-like” structure is formed (Fig. 1.10). Once the saw damage is etched away, the textured surface starts to flatten again for prolonged processing times, as sharp edges are rounded. Four to five micrometer removal of Si per side is normally enough to remove the saw damage and obtain a low reflectivity.8 The etching reaction takes place in two steps, an oxidation

8 Note that the maximum depth of saw damage can be up to around 10 μm, but as predominantly the damaged areas are attacked, less overall removal of Si is needed. State-of-the-Art Industrial Crystalline Silicon Solar Cells 21

3Si + 4HNO3 ! 3SiO2 +3H2O + 4NO (1.27) followed by etching of the SiO2

3SiO2 + 18HF ! 3H2SiF6 +6H2O: (1.28)

Afterward, the thin porous Si layer at the surface is etched off in (cold) KOH. The remaining reflectivity is significantly higher than for random pyramids, therefore it is not used for mono-Si (Fig. 1.11). Acidic texturing can be done elegantly inline, as texturing time is in the range of only around 2 min (depending on temperature) (Hauser et al., 2004; Neuhaus and Mu¨nzer, 2007).

Figure 1.10 SEM images of textured c-Si surfaces for mono-Si using KOH/IPA (left) and mc-Si using an acidic texture solution (right).

0.7 Bare Si 0.6 Si textured Bare Si AR coated 0.5 Si textured and AR coated 0.4

0.3 Reflectance 0.2

0.1

0.0 400 600 800 1000 1200 Wavelength (nm) Figure 1.11 Reflectivity of bare Si and alkaline-textured mono-Si with and without

SiNx:H ARC. 22 Giso Hahn and Sebastian Joos

After texturing, the wafers are thoroughly cleaned as the next step is the diffusion taking place at high temperatures. Impurities present on the wafer surface could diffuse in to the wafer causing recombination and therefore lowering τb. Cleaning normally consists of rinsing in deionized (DI) water, cleaning in HCl, DI water rinsing, etching in HF to form a hydrophobic surface, followed by a short dip in DI water and drying.

3.2. Phosphorus diffusion In this step, the heart of the solar cell, the pn-junction, is formed. The two most common ways to form the P-doped emitter will be described in the following. In most cases, the in-diffusion of P into c-Si takes place in a quartz tube furnace. N2 is directed into a bottle (bubbler) containing POCl3, which is liquid at room temperature. POCl3 molecules are transported with the N2 flow into the quartz tube, where the wafers are located in quartz containers (boats) with spacing between the wafers at temperatures around 800–850 C. O2 is added and on the wafer surface P2O5 is formed according to

4POCl3 +3O2 ! 2P2O5 + 6Cl2, (1.29) where the formed Cl2 provides an additional cleaning effect on the wafer surface. The O2 flow also oxidizes the Si surface, and the resulting SiO2 layer together with the P2O5 forms the so-called phosphor silicate glass (PSG) layer acting as the diffusion source. P diffuses into c-Si and the diffusion coefficient depends strongly on doping concentration, as the level of EF determines the amount of vacancies present in the material. For P concentrations well above 1019 per cm3 at diffusion temperature, a different diffusivity is observed due to the existence of double negatively charged vacancies in large amounts, forming a mobile quasi-particle with ionized P+. For P concentrations below 1019 per cm3, single negatively charged vacancies dominate the diffusion mechanism. This results in the characteristic kink-and-tail shaped profile of P diffusion in c-Si whereby the tail is formed due to the “normal” diffusion mechanism involving vacancies (Fair and Tsai, 1977). During diffusion, time, temperature, and gas flows have an influence on the diffusion profile formed. To increase (double) the throughput, wafers are often loaded in the quartz boats back-to-back. As an alternative to quartz tube POCl3 diffusion, a liquid P-containing layer can be deposited on the wafer surface (mainly diluted H3PO4), e.g., by spraying. Wafers then move horizontally through a conveyor belt firing State-of-the-Art Industrial Crystalline Silicon Solar Cells 23 furnace. As for this inline technique, the time allowed for diffusion is limited due to throughput and length of the furnace possible, diffusion temperatures are normally higher than for POCl3 diffusion, resulting in steeper P doping profiles. Higher doping concentrations normally result in lower emitter quality and more Auger recombination, increasing the saturation current density contribution of the emitter j0e. In addition, surface passivation is influenced by doping concentration with better passivation quality possible for lowly doped surfaces (Cuevas et al., 1996). Only P atoms on Si lattice sites are electrically active dopants. The surface concentration of P in c-Si for an unlimited source is given by the solubility limit in the range of 3–61020 per cm3 between 800 and 900 C with higher values for higher temperatures (Trumbore, 1960). Apart from the electrically active P atoms, interstitial P or P-containing clusters can form, increasing the amount of P present in Si especially close to the surface (Fig. 1.12)(Bentzen et al., 2006a). The almost flat shape of the P doping profile with electrically active P concentration above 1020 per cm3 is also referred to as “dead layer,” as this layer is highly recombination active. Although the high surface concentration of P close to the surface is limiting the electronic quality, it seems to be needed for the formation of a good low resistivity contact with the front metal Ag paste during the firing step. An important parameter of the emitter is its conductivity, as charge carriers have to flow laterally toward the collecting finger grid. As the emitter is a very thin layer (usually well below 1 μm thick), a sheet resistivity is defined for a uniformly doped layer as

Figure 1.12 P profiles of identical P diffusions in c-Si measured by ECV (electrically active concentration) and SIMS (total concentration). The solubility limit at diffusion temperature according to Bentzen et al. (2006b) and Solmi et al. (1996) is also indicated. 24 Giso Hahn and Sebastian Joos

ρ R ¼ sheet d (1.30) where ρ is the resistivity of the layer with thickness d. For non-uniformly doped layers as is the case for a diffusion profile, sheet resistivity calculates according to 1 Rsheet ¼ , (1.31) Ð d 1 dz 0 ρðÞz with depth z. Typical values of Rsheet for P diffusions are 50–100 Ω/sq. Note that carrier mobility is a function of doping concentration, therefore also conductivity is a function of doping density.

3.2.1 Phosphorus diffusion gettering of impurities

During P diffusion, SiO2 is formed at the wafer surface. For the formation of SiO2 Si, atoms have to leave their lattice sites, and a flux of Si interstitials Sii is generated. In addition, diffusing P atoms change position with a Si atom on lattice site (Sis), further increasing the concentration of Sii. These Si interstitials themselves can change position with an impurity atom (M) located at lattice sites via a “kick-out reaction”

Sii +Ms ! Sis +Mi: (1.32) In interstitial form, the impurity atom is mobile and can diffuse through the crystal toward a location with higher solubility. Such a location of high solubility can, e.g., be the highly doped region containing precipitated P close to the surface or the PSG (Bentzen et al., 2006a). Such regions are called getter sinks. Due to the presence of regions with high solubility for impurity atoms, a concentration gradient is formed toward the sink, resulting in more and more impurity atoms moving toward the sink, leaving a cleaner region behind. This self-cleaning process is called gettering of impurities. Depending on the location of the getter sink we distinguish between external gettering (e.g., at the crystal surface, where the impurities can be removed) or internal gettering at extended defects in the Si bulk (e.g., at grain boundaries, dislocations, or precipitates). The gettering process in general can be divided into three phases: (1) freeing the impurity from its (bonded) position by supplying an activation energy, (2) diffusion of the impurity in the wafer, (3) capture of the impurity at the gettering location. Depending on the specific mechanism present, it is distinguished between State-of-the-Art Industrial Crystalline Silicon Solar Cells 25 relaxation induced, segregation induced, and injection induced gettering, where also combinations of these different mechanisms are possible (Kang and Schroter,€ 1989; Seibt and Kveder, 2012). It could be shown that for back-to-back diffusion the positive gettering effect is less pronounced, as only one effective external getter sink is available (Schneider et al., 2005). The same can be assumed for inline diffusion, as the diffusion source is applied only on one side.9

3.3. Edge isolation After P diffusion, the PSG layer is etched in dilute HF and impurities gettered toward the PSG are removed. As the highly doped emitter is still present around the wafer edges and at the rear side, it will cause a short circuiting of the diode. It either has to be removed, or the connection between front and rear side has to be inter- rupted later in the process (e.g., by laser scribing; Emanuel et al., 2001). An elegant way to remove the rear side emitter is by inline etching of the rear side with the wafer floating on the etch solution containing H2O, HF, HNO3, and H2SO4. As the PSG etch can be performed inline as well, both steps can be combined in a single inline wet bench (Delahaye et al., 2004; Melnyk et al., 2005). Thereby care has to be taken that the emitter on the front side is not attacked by the etch solution or the atmosphere containing reactive species.

3.4. SiNx:H deposition To further minimize reflection losses at the front side, an ARC based on 10 SiNx:H is deposited on the front. For normal incidence of photons, destructive interference is reached if the thickness d of the ARC is λ d ¼ , (1.33) 4nSiN with wavelength λ and refractive index of the SiNx:H layer nSiN. Fresnel’s equations predict a zero reflectivity if pffiffiffiffiffiffiffiffiffiffiffi nSiN ¼ nairnSi, (1.34)

9 For back-to-back POCl3 diffusion and inline diffusion, a weak P diffusion is also observed at the rear side, by transport of P atoms via the gas phase. 10 The notation SiNx:H is used, as the silicon nitride layers are not of stoichiometric composition and contain significant amounts of H. 26 Giso Hahn and Sebastian Joos

with nair and nSi refractive indices of air and c-Si being the materials above and below the ARC, respectively. With nSi(600 nm)¼4, on cell level the optimum refractive index would be nSiN ¼2. As the cells are later encapsulated into modules using ethylene vinyl acetate (EVA) with refractive index around 1.5, a slightly higher nSiN would be optimum for solar cells in module application. With Eqs. (1.33) and (1.34) typical ARC thicknesses are 75–80 nm for nSiN ¼2, if reflectivity of photons with λ¼600–650 nm (the maximum in photon flux of the AM1.5g spectrum) should be minimized. Apart from the ARC effect, SiNx:H is also suited for surface passivation ++ of the n P-doped emitter. Defects located close to the c-Si/SiNx:H inter- 12 face in the SiNx:H layer provide fixed positive charges in the order of 10 – 1013 per cm2 (Aberle, 1999; Lamers et al., 2012). The minority charge carriers (holes in n-type material) are repelled from the surface due to Coulomb repulsion while majority carriers are attracted, and therefore recombination is lowered.11 This surface passivation mechanism is referred to as field effect passivation in contrast to chemical passivation, where the reconstruction of chemical bonds lowers the density of energy levels in the bandgap. Chemical passivation is also present for SiNx:H layers, but remaining defect densities are usually higher than for SiO2 layers, which in turn have a lower density of fixed charges. If the SiNx:H layer contains significant amounts of H, this H can be released during the firing step and diffuse into the c-Si bulk (Hahn et al., 2004; Jiang et al., 2003). Here, H can passivate bulk defects and drastically improve material quality (Duerinckx and Szlufcik, 2002; Hahn et al., 2010). This is a crucial step especially for mc-Si material with high defect densities. The most common technology to deposit SiNx:H layers is plasma- enhanced chemical vapor deposition (PECVD). There are direct and remote plasma techniques available. A direct plasma system usually operates at low frequency of 40 kHz. The wafer forming one electrode is in contact with the plasma, and accelerated ions can bombard the wafer surface leading to a certain surface damage.12 These systems usually operate in batch mode. In remote plasma systems, operating usually at high frequencies around 13.56 MHz, the plasma is spatially separated from the wafer, and a linear plasma source is used with microwaves supplying the excitation. Remote plasma systems are usually operating in inline geometries with the wafers

11 Note that recombination rate is dependent on np in Eqs. (1.10), (1.12), and (1.14). 12 This surface damage is not necessarily negative, as it can provide a reservoir for H. State-of-the-Art Industrial Crystalline Silicon Solar Cells 27 lying on trays being transported through the reactor underneath the linear plasma source. For both techniques, pressure is around 0.1–1 mbar and deposition temperature is between 300 and 450 C, depending on technology. Precursor gases used are SiH4 and NH3, and the SiH4/NH3 ratio determines the stoichiometry and therefore refractive index and absorption coefficient of the resulting SiNx:H. Higher SiH4/NH3 ratios (Si-rich layers) lead to higher refractive indices and higher absorption (Nagel et al., 1998). As absorption in the ARC is unwanted, usually a compromise between optimum nSiN (2.3–2.4 for module application) and low absorption is found, with nSiN(600 nm)¼2.0–2.1. Before PECVD SiNx:H with high throughput became available for PV at the end of the 1990s, TiO2 was often used as ARC for industrial c-Si cells. Compared to SiNx:H, a higher refractive index without significant absorption was possible, but TiO2 layers showed poor surface passivation qualities and contained no H needed for bulk passivation. The method used for deposition was atmospheric pressure chemical vapor deposition (APCVD).13

3.5. Metallization via screen-printing Screen-printing of metal pastes for PV application is a very robust method already introduced in 1975 (Ralph, 1975) and can be used as an inline process.14 A conveyor belt transports the wafer onto the printing chuck. A screen consisting of a mesh of wires partly covered with an emulsion is the mask for the metallization process. Metal paste is printed through the openings in the emulsion through the mesh of wires onto the wafer lying under the screen. The screen is positioned on top of the wafer with a well-defined distance between screen and wafer (the snap-off distance). The paste is placed on top of the screen and a squeegee moving horizontally without pressure on the screen fills the openings of the mesh uniformly with paste. In the next horizontal movement of the squeegee over the screen, it is pressed onto the screen with a defined pressure, pressing the screen locally against the wafer surface and pushing the paste from the filled areas onto the wafer surface. The screen snaps off from the wafer after the passing of the squeegee because of the screen tension (Fig. 1.13). After printing, the wafer is transported into a drying furnace for evaporation of the volatile ingredients of the paste at temperatures of around 150–200 C to avoid smearing of the paste when it is flipped over to metallize the other surface. Due to the

13 See Richards (2004) for a review on TiO2 and other dielectrics for use of ARC. 14 See also Holmes and Loasby (1976) and Neuhaus and Mu¨nzer (2007) for more details. 28 Giso Hahn and Sebastian Joos

Squeegee

Frame Metal paste Open mesh Closed mesh

Vacuum chuck Solar Cell

Figure 1.13 Screen-printing of metal paste. tension put on the screen during every printing step, the lifetime of the screens is limited to several thousand printing steps as they wear out with time.

3.5.1 Front side metallization For front side metallization, the following criteria have to be met by the paste used: (1) low contact resistance to c-Si, (2) low specific resistance in the printed structure, (3) no junction shunting, (4) good aspect ratio (height to width ratio) of the fingers, (5) good adhesion to c-Si, (6) opening of SiNx:H layer provided, and (7) solderability for cell interconnection in the module. Ag containing paste (70–80%weight) is used, as Ag is highly conductive and therefore allows for good conductivity in the printed metal fingers. Additional components are glass frits containing PbO, B2O3, and SiO2 (1–10%weight), which are responsible for locally dissolving the SiNx:H layer State-of-the-Art Industrial Crystalline Silicon Solar Cells 29 as well as for a good adhesion. Also present are organic binders (15–30%weight), influencing the rheology of the paste which has to be low enough to ensure that a continuous finger is formed and high enough to keep a high aspect ratio (Neuhaus and Mu¨nzer, 2007). During the firing step at temperatures around 800 C, the electrical contact between Ag and c-Si is established. Early detailed studies for Ag front contact formation by Ballif et al. (2002, 2003) and Schubert et al. (Schubert, 2006; Schubert et al., 2002, 2004) led to the following picture (Fig. 1.14). Below 600 C organic components burn out (A). At higher temperature, the contact is formed as first the PbO melts, wets, and etches the SiNx:H layer (B). The Ag particles with sizes of several μm sinter together and form a conductive film. Then a redox reaction between PbO and Si forms Pb (C). The liquid Pb starts to melt Ag (D), and the Ag/Pb melt reacts with Si, etching inverted pyramids locally into the c-Si surface (E). On cooling down Ag recrystallizes on (111)-Si planes, forming isolated contact points to the emitter (F). The recrystallized Ag points at the c-Si surface can either be in direct contact with the sintered Ag layer, or the glass layer iso- lates them from each other. If there is no direct contact established, contact resistance depends strongly on the thickness of the glass layer formed in between. Thin layers can be tunneled through, with small (nanoscale) metal precipitates of Ag and/or Pb/Bi providing additional hopping sites for electrons. The thickness of the isolating glass layer is a very crucial parameter for achieving low contact resistance in case of no direct connection between Ag crystallites and sintered Ag layer. Therefore, firing parameters are very important, with too high peak temperature resulting in thicker glass layers and too low peak temperature resulting in not completely opened SiNx:H layers (Schubert, 2006). As the contact is not formed everywhere underneath the Ag metallization printed onto the wafer, the contact resistance is significantly higher than + for contacts established, e.g., via evaporation of Ag directly on c-Si n emitters. Typical values for contact resistance of screen-printed Ag on c-Si are 1–10 mΩ cm2 (Schubert, 2006), while values for evaporated contacts are in the range of 100–200 μΩ cm2 (Fischer, 1994).

3.5.2 Rear side metallization For rear side metallization, Al containing pastes containing Al powder, glass frit, organic binders, and solvents are used. The lower conductivity of Al compared to, e.g., expensive Ag does not play a role as long as the contact is formed on the full area rear side. Another very important advantage of Al is 30 Giso Hahn and Sebastian Joos

ABC

Silver Silver Silver

- g - g

DEF

Silver Silver Silver

Figure 1.14 Schematic contact formation for Ag screen-printing on n++ emitters. After Schubert (2006). the fact that Al is an acceptor in c-Si and can form a good ohmic rear contact in combination with a highly p+-doped layer, the BSF (Mandelkorn and Lamneck, 1972). The BSF is formed by alloying during the firing step after drying of the paste (Lolgen,€ 1995). A description of the formation process is given in Huster (2005). The Al/Si phase diagram depicted in Fig. 1.15 shows the composition of the Al/Si melt dependent on melt temperature. Upon heating up during the firing step, Al starts to melt at 660 C. The Al2O3 shells around the Al spheres stay in shape, but liquid Al can penetrate through the oxide shell locally and gets in contact with the c-Si surface and other Al particles. Si is dissolved into the Al melt at that temperature until according to the phase diagram the melt contains around 17% Si. As the volume in the stable oxide shells stays constant, the amount of Al leaving the shell covering the c-Si surface corre- sponds to the volume fraction of Si entering the shells to form the correct concentration according to the phase diagram. During further increase of temperature, more and more Si is dissolved in the melt. Assuming a peak firing temperature of 800 C, the melt contains around 27% of Si. During cooling down from peak temperature, a part of the Si has to leave the melt according to the phase diagram. This Si recrystallizes at the c-Si/melt interface. During recrystallization, a small amount of Al is incorporated into the recrystallized Si lattice according to the solid solubility of Al in Si at that State-of-the-Art Industrial Crystalline Silicon Solar Cells 31

L (T) 1414 °C 1400 Liquid Al–Si-phase 1200

1000 Solubility of Al in Si 800 T Liquidus curve L (T) T 4 peak Solubility of Si in Al T 3 2 T 5 Temperature (°C) 600 577 ± 1 °C T 6 eut 12.2% (Al) + (Si) 400

0 20 40 99.98 100

AI %at silicon Si

Aluminium paste

Al solid Al liquid Si wafer AlSi liquid Si BSF T Њ T Њ T Њ T = 700 ЊC T <577 ЊC Paste dried 2 = 660 C 3 = 700 C 4 peak = 825 C 5 6 AlSi solid melting of Al start of alloying 1. 2. 3. 4. 5. 6. Figure 1.15 Phase diagram of Al/Si (top) with characteristic stages during rear side contact firing (bottom). Al/Si data from Krause et al. (2011) and solubility data from Murray and McAlister (1984) and Yoshikawa and Morita (2003). After Huster (2005). temperature. This accounts for the p+ doping of the recrystallized Si layer in the range of 1018–1019 per cm3, the BSF. The recrystallization stops when the temperature of the melt reaches the eutectic point at 577 C. The solidified Al/Si melt then has eutectic composition with 12% Si. This holds true for both the solidified melt in the oxide shells as well as for the film directly on top of the recrystallized c-Si (BSF). As solubility of Al in Si decreases from peak temperature toward 577 C, the BSF contains a doping gradient with higher Al concentration at the p/p+ interface. The thickness of the resulting BSF in dependence of the amount of Al deposited can be calculated according to g F E d ¼ Al BSF ρ F E , (1.35) Si 100 100 2 with gAl the amount of Al deposited (in g/cm ), ρSi the density of Si, F the fraction of Si in the Al/Si alloy at peak temperature and E at eutectic temperature in percent (del Alamo et al., 1981). 32 Giso Hahn and Sebastian Joos

The p/p+ interface forms a high/low junction, and similar to the pn-junction the doping gradient is the source for an electric field causing slight band bending close to the rear metal contact (see Fig. 1.5). This results in repelling of electrons (minority carriers) and therefore less recombination at the semiconductor/metal interface in analogy to surface passivation via the field effect. But as the BSF region is highly doped, its electronic quality (minority carrier lifetime and diffusion length) is quite low. The effect of BSF parameters on surface passivation can be calculated via (Godlewski et al., 1973) s L + d b + tanh BSF N D + D + L + s ¼ A , (1.36) eff N + L + s L + d A 1+ b tanh BSF D + L + with seff the effective back surface recombination velocity at the edge of the + + + quasi-neutral region of the base, NA , L , and D the acceptor concentration, the diffusion length, and the diffusion constant of minority carriers in the BSF, respectively, and sb the surface recombination at the physical back + surface. It can be seen that for low seff, e.g., a high L and a low ratio of NA/ + NA (high concentration gradient) are beneficial. Therefore, B can be mixed to the Al paste to increase p+ doping of the BSF, as B has a higher solubility in Si compared to Al (Rauer et al., 2013). As cells have to be interconnected for module integration and the Al rear side is not solderable, a small fraction of the c-Si rear side is metallized with Ag or Al/Ag paste pads or busbars. These regions have to be kept as small as possible to avoid unnecessary recombination, as in these regions no BSF is 15 formed and seff is significantly higher. On the other hand they have to be large enough to assure reliable soldering.

3.5.3 Co-firing step Contact formation is realized in a conveyor belt furnace with optical heating. The simultaneous contacting of emitter and base in the so-called co-firing step is a very critical part of the process, as it has to be optimized for several purposes. During this step, BSF and rear contact are formed, while in addition on the front side the Ag metallization has to be fired through the SiNx:H layer, and H has to be released from the SiNx:H layer into the c-Si bulk to passivate crystal defects. As optimum parameters for

15 Surface recombination velocity of a pure c-Si/metal contact is in the range of 106 cm/s. State-of-the-Art Industrial Crystalline Silicon Solar Cells 33

Figure 1.16 Typical firing profile. Data from Huster (2005). some of these steps are going in different direction of the parameter space, a compromise has to be found. Ag and Al paste as well as emitter profile and SiNx:H layer therefore have to be tuned to match their optimum firing parameters as good as possible and allow for a good overall end result. In Fig. 1.16, a typical firing profile is shown. (1) the first plateau visible refers to the temperature of 660 C, where Ag starts to melt (latent heat). The peak temperature of around 800 C (2) is kept for only a few seconds. During cooling down, another plateau at the eutectic point of 577 C can be seen (heat of crystallization) (3). If edge isolation was not already performed earlier in the process (by single side etching or other methods like, e.g., plasma edge isolation), laser edge isolation close to the cell’s edge at the front side is another option. Here, the emitter is locally removed by laser heating, and a groove is formed (Emanuel et al., 2001; Schneiderlochner€ et al., 2003). Hereby active cell area is lost, slightly compromising the current and therefore efficiency (Hauser et al., 2001).

3.6. Solar cell characterization After solar cell processing, the j–V characteristics of every cell are measured, with Voc, jsc and FF determining the conversion efficiency (Eq. 1.6). From these, the current density and voltage at mpp are determined and solar cells are classified into bins according to current density classes under mpp conditions to avoid current mismatch in the string of the module when cells are 34 Giso Hahn and Sebastian Joos connected in series. In addition, all cells are inspected visually to assure that only solar cells with the same color end up in a module to achieve homogeneous optical appearance. Also the reverse break through characteristic of the diode is checked by applying reverse bias of 12–15 V to avoid the formation of hot spots, which might cause dangerous local heating (hot spots) of the cell. This could happen when, e.g., one cell or part of it is shaded in the module and current is flowing in reverse direction through the string, leading to destruction of the cell. j–V characteristics are normally measured using a halogen flash lamp that can provide a constant power of 1000 W/m2 for around 50 ms, with the light intensity checked by a monitor cell. The actual temperature of the solar cell is measured and the voltage is corrected to the temperature of STC, 25 C. As for STC, the spectrum should be AM1.5g, which is never achieved exactly by the flasher system, certified reference solar cells have to be used for calibration to minimize the effect of spectral mismatch. The calibration cell should exhibit an external quantum efficiency (EQE) very similar to the cell to be measured to avoid errors, especially when introducing new cell designs (Herguth et al., 2011). Contacting and measurement are realized using the four point probe setup, with several probes for each busbar simulating the situation after tabbing of the cells for interconnection in the module. Solar cell parameters are normally determined under STC conditions, but solar modules under realistic operation conditions in the field might operate at significantly different temperatures. Depending on the location of operation, solar cell temperatures well above 25 C are reached, reducing the efficiency mainly according to the temperature behavior of the voltage. This should be kept in mind when optimizing not only the solar cell, but the complete PV system for an optimized energy output.

4. RECENT DEVELOPMENTS ON SOLAR CELL FRONT SIDE 4.1. Wafer sawing The standard technology for sawing of wafers is still slurry-based cutting of wafers using a steel wire as described in the previous section. As the wire has a diameter of around 120 μm, resulting in kerf loss of around 140 μm, and wafer thickness is typically 180 μm, around 40–50% of Si is lost during this step. The slurry used for cutting and cooling normally contains SiC particles for cutting and other additives contaminating the Si kerf. Therefore, the contaminated Si kerf cannot be easily recycled and is lost. State-of-the-Art Industrial Crystalline Silicon Solar Cells 35

A lot of effort went into the development of alternative sawing processes, with diamond wire sawing as the most promising alternative. Here, the wire is coated with small particles (diamond-plated wire), and the wire itself is the cutting source. No additional slurry is needed, and water is used as a cooling agent. In this way, Si kerf can in principle be separated and recycled. After diamond wire cutting, the surface damage of the wafer is different compared to slurry-based cutting (Buchwald et al., 2013). The damage is generally not as deep as for slurry-based cutting, but is less homogeneous. In addition, the surface exhibits ripples, which might have an influence on mechanical stability and processing steps following later during solar cell processing.

4.2. Alkaline wafer texturing Wafer texturing based on alkaline solutions with IPA as an additive shows limitations in throughput, cost, and reliability. As IPA has a boiling point of 82 C, texture bath temperature is limited to around 80 C. This causes throughput issues, as the underlying chemical texture reaction increases with temperature. In addition, IPA constantly evaporates during texturing. Therefore, it has to be replenished to assure the correct composition of the texture bath. A lot of research was carried out to find alternatives to KOH/IPA texture solutions with industrial relevance. Some of them try to minimize the consumption of IPA (see Basu et al., 2013 for an overview), others try to avoid it completely. IPA free texturing recipes/processes (see, e.g., Moynihan et al., 2010; Ximello et al., 2009) using additives/alcohols with higher molecular weight have the advantage that higher texture bath temperatures can be used, speeding up the texturing process and leading to more homogeneous texturing results. This is of high interest especially for wafers sawn with the new diamond wire technique or treated with different washing/cleaning solutions prior to texturing, as standard KOH/IPA texture solutions often fail to texture these wafers reliably. An alternative to avoiding IPA completely might be the application of a closed chamber texture bath, with the texturing process being sped up by application of vacuum pulses and automatic recycling of evaporated IPA (Ximello et al., 2011).

4.3. Front contact metallization Although screen-printing is still the most reliable and cost-effective way of forming contacts in solar cell mass production, the standard Ag screen- printing technology has several drawbacks and challenges. 36 Giso Hahn and Sebastian Joos

• Poor aspect ratio resulting in shadowing losses • Saturation current density is drastically increased at the metal/semiconductor interface • Relatively high contact resistivity between metal and semiconductor • Series resistance losses increase when solar cell is interconnected in the module (cell to module losses) In the following, new developments are presented, addressing the issues mentioned above.

4.3.1 Double print The aspect ratio of screen-printed Ag fingers can be improved by using a screen with narrower openings. This increases the risk of finger interrup- tions, therefore a second printing step on top of the already existing metal grid finger line provides continuous metallization and increased finger height, improving the aspect ratio (Galiazzo et al., 2009). For this approach, an additional print step is needed.

4.3.2 Dual print The front metal grid normally consists of the fingers and three collecting busbars running perpendicular to the fingers. The purpose of the fingers is to collect the current in the emitter, while the main purpose of the busbars is to collect the current from the fingers. In standard screen-printing, the front grid is printed in one step, with the glass frits present in the paste etching the SiNx:H layer. Therefore, also the area under the busbars has direct contact to the emitter, as the SiNx:H layer is etched away. In principle, this additional area is not needed for contacting, but increases the saturation current density. For an estimation, 1 mm wide busbars and 70 μm wide fingers with a finger distance of 2 mm lead to a metal coverage of 5.3% for a 156156 cm2 solar cell, with busbars being responsible for 1.9%. If printing of fingers and busbars is separated in two separate printing steps, for busbars a fritless Ag paste can be used, reducing recombination, because underneath busbars the passivating SiNx:H layer is still present. But here again an additional printing step is needed.

4.3.3 Stencil printing For stencil printing (Fig. 1.17) not an elastic screen with a mesh, but a fixed metal plate with openings is used (de Moor et al., 2012). Advantages are that the stencil allows for narrower fingers and a better aspect ratio while showing less wear-out than a screen, which can only be used for a limited amount of State-of-the-Art Industrial Crystalline Silicon Solar Cells 37

Squeegee

Frame Metal paste Opening Stencil

Vacuum chuck Solar cell

Figure 1.17 Schematic for stencil printing. prints (usually several thousand) before it has to be exchanged. Stencil printing can, e.g., be combined with dual print for achieving narrow fingers. Disadvantages are higher costs for the stencil and the fact that not all geometric patterns can be printed in one step. For a comparison of different new printing options see, e.g., Hannebauer et al. (2013).

4.3.4 Dispensing Another alternative to standard Ag screen-printing for deposition of Ag on the wafer surface is dispensing (Hanoka, 1989; Hanoka and Danielson, 1991). Here, the paste is pushed through ceramic nozzles, resulting in significantly less spreading of the fingers on the wafer, resulting again in better aspect ratios (Specht et al., 2010). For this application, rheology of the paste has to be optimized (Pospischil et al., 2011, 2012), but no extra step is needed, making this technique very interesting for industrial application 38 Giso Hahn and Sebastian Joos

(Beutel et al., 2014). A similar extrusion technique has already been used for metallization of EFG solar cells in the past (Tobias et al., 2003), and as the technique is putting much less pressure on the wafer than screen-printing, it is especially well suited for thinner or more fragile wafers.

4.3.5 Paste development The contact formation between emitter and Ag paste is highly dependent on paste formulation and properties of the Si surface. The goal is to establish good contact with low contact resistivity on emitters with low P surface concentration in combination with high conductivity of the Ag finger. This would enable low series resistance (high FF) and low recombination losses (good surface passivation with high Voc and jsc). In reality, there is a trade-off between low surface doping concentration and high contact resistivity. Nevertheless, over the past years Ag paste formulation was changed by the suppliers, allowing for higher sheet resistivities16 being reliably contacted (Cabrera et al., 2013; Carrol, 2013). The mechanism of contact formation of these newly developed pastes is still under investigation, but at least for some pastes smaller Ag particles in combination with a tunneling (or hopping) mechanism of electrons might explain the lowered contact resistivity on lower doped Si wafer surfaces (Cabrera et al., 2011; Li et al., 2011). In addition, the thickness of the glass frit layer between Si and Ag particles in the paste plays an important role, with very thin layers showing best results. Standard sheet resistance in industry today for homogeneous emitters is around 70–80 Ω/sq.

4.3.6 Seed-and-plate The application of Ag metal paste for front contact formation is always a compromise between low contact resistivity, good line conductivity, and high aspect ratio (low shadowing). An alternative addressing these issues is the seed-and-plate approach, where first a seed layer is formed, establishing low contact resistance, but not necessarily showing high conductivity. In a second step, the contact is thickened by plating of highly conductive metal. If in the first step a narrow seed layer is formed, high aspect ratios can be reached after plating. Several methods can be used to form the seed layer. The first industrial concept related to this approach was the laser grooved buried contact Saturn solar cell concept from BP Solar (Mason et al., 2004).

16 Rsheet is not a very good characteristic for the quality of the emitter, as emitter quality is more closely linked to j0e. Nevertheless, sheet resistivity can be measured in an easy way, therefore it is often referred to. State-of-the-Art Industrial Crystalline Silicon Solar Cells 39

Here, electro-less plating of a Ni layer is used, and after sintering forming nickel silicide low contact resistivity is reached. The following Cu plating step provides excellent line conductivity (Wenham and Green, 1985). This concept was used in conjunction with a selective emitter structure based on two P diffusions and a laser step to form grooves for the fingers to be plated into. As this concept needed several extra steps compared to the standard process, it was abandoned. As an alternative to Ni plating, an adapted Ag paste can be deposited and used as a seed layer for the following Ag plating step. The Ag seed layer can be very thin and should form narrow fingers. Deposition can, e.g., be performed via aerosol printing. In this way, the Ag paste formulation can be optimized toward low contact resistivity and low shadowing. This can be done by adapting the glass frit content (see, e.g., Horteis,€ 2009 for more details). The Pluto solar cell concept of Suntech uses also plating for front side contact formation in combination with a selective emitter (Shi et al., 2009).

4.3.7 Multi-busbar approaches For optimized power output on cell level, three busbars are normally used on the screen-printed Ag front metallization grid for cells of 156156 mm2 size. This is the optimum amount of busbars of standard dimensions17 in case that j–V measurement is performed on solar cell level via three contacting bars each holding a multitude of current pins and a central voltage pin each to minimize Rs losses of the busbars. After interconnection of the individual 18 solar cells using ribbons, the Rs loss of the string is significantly different from the situation of the individual solar cells. Total power loss in the string

2 2 Ptot ¼ Rs,totI ¼ Rs,totðÞjA (1.37) is dependent on total series resistance including the ribbon interconnectors and the current I squared (depending on the cell area A). Therefore, under module conditions the optimum amount of busbars is different, as Rs,tot is increased due to the contribution of the connecting ribbons. Several approaches have been developed in recent years to address this issue. As adding more and more screen-printed busbars of standard size would increase shading losses, only fingers are applied by, e.g., using the standard

17 On top of the busbars Sn coated Cu ribbons are soldered for interconnection of the cells in a string. Typical dimensions of the ribbons are widths of 1.2–1.5 mm with a thickness of 150–250 μm. 18 Referred to as module conditions 40 Giso Hahn and Sebastian Joos screen-printing technique. Then a multitude of coated Cu wires are substituting the screen-printed busbars. The wires could either be connected to a foil (Schneider et al., 2006) pressed and soldered onto the cell surface, or could only be soldered to the fingers at the crossings between wire and finger (Braun et al., 2013a). Choosing the correct amount of wires with given diameter, higher efficiencies under module conditions can be reached, because of the new trade-off between shadowing losses and Rs (Braun et al., 2012). In addition, with negligible contribution of finger line resistance to Rs using more busbars reducing effective finger length, a significant further reduction of Ag consumption is possible, apart from saving the Ag used for the screen-printed busbars (Braun et al., 2013a, 2014). This approach is also well suited to be combined with a seed-and-plate approach, further reducing the Ag consumption and leading to higher efficiencies (Braun, 2014; Braun et al., 2013b). It is also an elegant way to interconnect solar cells fabricated with the heterojunction with intrinsic thin layer (HIT) concept (Papet et al., 2011).

5. ADVANCED EMITTER FORMATION

The pn-junction can be described as the heart of a c-Si solar cell, as it is responsible for the separation of charge carriers. For emitter formation within the standard solar cell process with a B-doped p-type base and a P-doped n-type emitter several factors have to be considered, the most important ones are listed in the following. 19 • j0e should be low enough to allow for high Voc values (see Eq. 1.24) • Rsheet should be low enough to provide good conductivity for lateral carrier transport toward the grid fingers • Surface concentration of P atoms should be high enough to allow contacting of the emitter via Ag paste • Surface concentration of P atoms and dead layer thickness should be low enough to minimize Auger recombination and increase blue response20 An optimization of emitter quality is therefore not a straightforward task, as the requirements listed above point in different directions. j0e is normally measured using symmetrical lifetime samples with the emitter and the

19 j0e is part of the saturation current density j0 (together with contributions from the bulk and the SCR), j0e is influenced by the surface passivation layer on top of the emitter. 20 Spectral response and IQE of short wavelength photons with low penetration depth is reduced when the very highly (concentration >1020 per cm3) doped layer at the emitter surface is too thick. State-of-the-Art Industrial Crystalline Silicon Solar Cells 41 covering surface passivation layer present on both sides. Kane and Swanson (1985) showed that j0e can be derived from the measured lifetime

1 1 2s 1 2j ðÞNA + Δn ¼ + ¼ + 0e : (1.38) τ τ d τ dqn2 eff b b i Using lowly doped high quality material (FZ Si) and high injection conditions (ΔnNA for p-type material) 1/τb and NA in Eq. (1.38) can be neglected, and j0e can be easily derived. j0e is influenced by both Auger recombination in the emitter as well as surface passivation (e.g., the SiNx:H ARC). There is a trend toward lower j0e with lower P surface concentration (Book et al., 2010) and lower Rsheet (Book et al., 2011), but Rsheet alone is not a very reliable parameter for evaluation of emitter quality, as will be shown in the following.

5.1. Improvement of homogeneous emitters For not too high P concentration (<1020 per cm3), the electrically active21 P concentration equals the total P concentration. This might not be the case for concentrations >1020 per cm3, because close to the solubility limit of P in Si clusters containing P atoms can form. In this case, the electrically active P concentration can be significantly smaller compared to the total P concentration (up to one order of magnitude). Electrically inactive P can form defects, increasing charge carrier recombination in this region. Therefore, from a recombination point of view a high quality emitter contains all P atoms in electrically active form. Nevertheless, a certain amount of electrically inactive P might be needed to obtain a low contact resistance for Ag screen-printed front contacts. Recently, a lot of effort went into a better understanding of the emitter formation process and toward a better understanding of the impact of POCl3 diffusion parameters on emitter quality for screen-printing metallization. It could be shown that the ratio of POCl3–N2 to O2 flow (Fig. 1.18) plays a crucial role for the thickness of the formed PSG (Dastgheib-Shirazi et al., 2012). P diffusion is a process known from microelectronics for quite some time, and very good tools for simulation of P diffusion in Si exist. But the diffusion coefficient of P in the diffusion source (PSG) at diffusion temperatures is still not completely understood. Detailed studies of the P concentration profiles in the PSG (Steyer et al., 2012) showed that the PSG is not a homogeneous layer, but consists of different regions, indicating

21 P atom on substitutional lattice site, acting as a donor. 42 Giso Hahn and Sebastian Joos

N2 POCl3 –N2,O2,N2 O2,N2

N 2 N2

Heat up Temperature Stabilization Diffusion Drive in Cool down

Time

Figure 1.18 Schematic visualization of the different stages during POCl3 diffusion with varying gas flows. that part of the PSG is liquid at diffusion temperature, whereas other regions are solid (Micard et al., 2012). Models for simulation of P diffusion (Wagner et al., 2011) can be improved with these findings. Implementing these results into solar cell processing via a design of experiment approach (Dastgheib-Shirazi et al., 2013a) led to significant improvement of Voc, jsc, and η of standard industrial Cz solar cells with full area Al-BSF. Efficiencies of 19.4% could be reached on a 55 Ω/sq emitter using standard single print screen-printing, while maintaining high FF with standard commercial Ag paste (Dastgheib-Shirazi et al., 2013b).

5.2. Selective emitters An elegant way to solve the problem of different needs for different emitter regions is a selective emitter structure.22 A selective emitter allows decoupling of the metallized and nonmetallized emitter areas. While the contacted area via screen-printing a high doping concentration at the surface and a deep emitter is beneficial because of the resulting lower contact resistance and the wide firing window, the nonmetallized areas need a lower doping level at the surface resulting in less (Auger) recombination and better surface passivation. High doping underneath the contacts can result in a

22 Part of the following discussion is based on a review on selective emitter technologies (Hahn et al., 2010). State-of-the-Art Industrial Crystalline Silicon Solar Cells 43 higher FF, and the lower doping in the nonmetallized areas results in better blue response with higher jsc and higher Voc values due to better surface passivation (and thinner or negligible dead layer). Selective emitters have been applied in lab-type processes in order to reach very high efficiencies for a long time (see e.g., Zhao et al., 1996). The first integration into an industrial-type process was via the buried contact approach (Wenham, 1993) which was commercialized by BP Solar. While higher efficiencies than for standard industrial-type screen-printed cells have been reached, process complexity was a drawback, as extra steps (e.g., laser groove formation, low pressure (LP) CVD of SiNx, second diffusion at high temperatures, Ni/Cu plating) meant extra costs, and a non- standard cell fabrication line layout was needed. In addition, the temperatures during the second P diffusion did not allow for a hydrogenation of bulk defects via a SiNx:H layer due to out-diffusion of H (apart from the fact that PECVD SiNx:H exhibits pin holes in contrast to the more dense LPCVD SiNx and can cause parasitic plating). This process was therefore not suited for the processing of mc-Si solar cells. From this experience, some conclusions can be drawn. For successful implementation of a selective emitter process into industrial mass production, several aspects have to be considered which form a wish list: • A minimum of extra steps • Possibility of implementation into existing cell lines • No yield losses (high stability and reliability) • Higher efficiencies (also for mc-Si) • Higher efficiency not only on cell, but also on module level. As a rule of thumb, efficiency should be increased by 0.2%abs for every extra step needed. Having in mind the points discussed above, several selective emitter technologies have been developed over the past years for the purpose of implementation in industrial mass production. Several of them will be presented, with the restriction to those which are already in production (or have at least been tested on pilot line level) and for which published academic information is available. The list therefore might not be complete, but is intended to serve as an overview of the various possibilities to realize a selective emitter structure. Further restrictions are the full Al-BSF which allows compatibility with existing cell technology and the possibility to use screen- printing for front side contact metallization (although some of the presented technologies develop their full potential with alternative front side metallization like plating). 44 Giso Hahn and Sebastian Joos

5.2.1 Doped Si inks Innovalight Inc. developed a technology based on highly doped Si nanoparticles which can be deposited onto the Si wafer surface via screen-printing prior to P diffusion (Antoniadis et al., 2010). Hereby, the ink is deposited only in the areas where the screen-printed front contact is located afterward. In the subsequent P diffusion step a lowly doped emitter is realized in the uncovered areas (80–100 Ω/sq) whereas the areas with the highly doped Si nanoparticles serve for contacting (30–50 Ω/sq). This technology adds only one additional step to the cell process prior to P diffusion (see Fig. 1.19).

5.2.2 Oxide mask process Centrotherm presented a selective emitter technology based on a masked P diffusion, where a thin structured SiO2 layer slows down the diffusion of P atoms from the surface into the Si bulk underneath the SiO2 (Esturo-Breton et al., 2009). Structuring of the SiO2 is done via laser ablation of the areas where the contacts are formed afterward. A wet chemical etching step removes the damage induced by the laser. The heavily doped region (300 μm wide) exhibits a sheet resistance of 45 Ω/sq and the masked area 110 Ω/sq. This technology offers a certain degree of freedom in emitter formation and uses technologies already established in PV.

5.2.3 Ion implantation process Varian introduced a new technology for selective emitter formation based on ion implantation through a mask which reduces the implanted dose in the areas between the contacts (Low et al., 2010). An annealing step in oxi- dizing ambient is carried out for crystal damage removal caused during

SE process Extra step(s) P-diffusion Extra step(s) P-glass etch

PECVD SiNx:H Screen-print front Screen-print rear Extra step(s) Co-firing Figure 1.19 Standard processing scheme for homogeneous emitter solar cells (right), with extra steps to be added for various selective emitter (SE) approaches. State-of-the-Art Industrial Crystalline Silicon Solar Cells 45

implantation and forms a thin SiO2 layer on the wafer surface acting as surface passivation. The process continues with SiNx:H deposition. Advantages of this approach are the dry processing for emitter formation, the lack of PSG formation (which normally has to be removed wet chemically) and of junction isolation, as the emitter is formed only on the front side. In addition, the amount of process steps is not increased.

5.2.4 Etch-back process University of Konstanz, Germany, developed an etch-back process which removes the dead layer of the heavily diffused regions after screen- or inkjet-printing of a mask covering the areas where the contacts are formed afterward (Haverkamp et al., 2008). The etch-back is performed via the formation of porous Si and allows for a very sensitive and controllable removal of the first tens of nm, as the porous Si formation is slowed down with increasing layer thickness almost independently of crystal orientation. This process adds only one new tool for masking, as the porous Si formation as well as etching of porous Si, PSG, and mask can be performed in the same wet bench already used for edge isolation, by adding more chemical baths. In addition, it uses only existing technologies and has been commercialized by Gebr. Schmid. As this approach of forming a selective emitter currently has the largest market share of the technologies described in this section (Gabor, 2012), some more information will be given. Figure 1.20 illustrates the principle

1021 2ϫ1020 cm−3 )

3 20

− 10 52 Ω/sq 17 Ω/sq 1019

1018

Calculated

P concentration (cm 17 10 sheet resistance: 118 Ω/sq 73 Ω/sq

0.0 0.2 0.4 0.6 0.8 1.0 Depth (µm)

Figure 1.20 Principle of emitter etch-back with removal of the highly doped dead layer and the possibility of tailoring the doping profile by, e.g., etching-back to the same Psurf 20 3 of 210 per cm resulting in different values for Rsheet. Data from Book et al. (2009). 46 Giso Hahn and Sebastian Joos

1021 R Ω Calculated sheet from all profiles: 106 /sq

Direct diffusion 1020 etch-back from 53 Ω/sq Ω

) 30 /sq

− 3 10 Ω/sq 1019 (cm D N

1018

1017 0.0 0.1 0.2 0.3 0.4 0.5 Depth (μm)

Figure 1.21 Different profiles with the same Rsheet after etch-back for different starting values (see Book, 2014). of the emitter etch-back. The dead layer of a heavily diffused emitter is etched-back until the desired phosphorous surface concentration Psurf is achieved. The result is a relatively deep P-profile with a low Psurf but still relatively low Rsheet, which cannot be reached by direct diffusion. In this way, Rsheet and Psurf can be decoupled to a certain degree (Fig. 1.21), and the emitter conductivity does not have to be increased, allowing the same front grid finger spacing without FF losses due to a higher Rs. Lowering of Psurf reduces the emitter saturation current density j0e. This can be seen in Fig. 1.22 where different directly diffused emitters are etched- back step by step. As a result, very low j0e values can be reached, independently of the starting Rsheet because j0e is mainly influenced by Psurf. As no high temperature steps exceeding the POCl3 diffusion temperature are involved, the etch-back selective emitter technology is also well suited for mc-Si. A similar increase in efficiency of up to 0.5%abs has been observed. The improvement by etch-back is more pronounced for inline emitter formation (see Section 3) as compared to POCl3 diffusion, reaching a similar quality as the homogeneous POCl3 emitter after etch-back (Hahn, 2010). This is due to the typically higher Psurf of inline emitters due to higher diffusion temperatures and shorter diffusion times compared to POCl3 emitters. Therefore, the blue response is normally lower for inline emitters after diffusion, and they can benefit more from reducing Psurf during the State-of-the-Art Industrial Crystalline Silicon Solar Cells 47 e j

Figure 1.22 Lowering of j0e with increasing etch-back (open symbols) of directly diffused POCl3 emitters (solid symbols). Similarly low values for j0e can be achieved by etching-back from different Rsheet starting values (emitter passivation by SiNx:H and firing). Data from Book (2014).

n 680 SiNx:H = 2.15 n = 2.0 675 SiNx:H

670 (mV)

oc 665

660 impl. V 655

650

40 50 60 70 80 90 100 R Ω sheet ( /sq)

Figure 1.23 Dependence of implied Voc on Rsheet for a POCl3 emitter etched-back from 40 Ω/sq. A fired PECVD SiNx:H layer with higher refractive index leads to higher implied Voc values, especially with increasing Rsheet (lower Psurf). Data from Haverkamp (2009). etch-back process. But the best performance is still reached for etched-back POCl3 emitters. The lower doped part of selective emitters is more sensitive to surface passivation than a homogeneous emitter, mainly due to lower Psurf. This effect is demonstrated in Fig. 1.23 by implied Voc values of an emitter etched-back from 40 Ω/sq. The emitter is covered with two different 48 Giso Hahn and Sebastian Joos

PECVD SiNx:H layers differing in refractive index. The SiNx:H with higher refractive index of nSiN ¼2.15 leads to higher implied Voc values, especially for higher Rsheet (lower Psurf). Therefore, a stack system of PECVD SiNx:H layers with a thin highly refractive and well passivating layer followed by a thick layer with standard refractive index can further increase surface passivation and cell performance (Dastgheib-Shirazi et al., 2009).

5.2.5 Laser doping via P-glass University of Stuttgart, Germany, introduced a laser-based selective emitter technology, where the P-glass present after a 110 Ω/sq P diffusion acts as P-source for the following laser process (Eisele et al., 2009; Roder€ et al., 2010). The laser with a special line-shaped beam profile melts the surface region in the areas for later front contact formation, and the recrystallized region is highly P doped without crystal defects. The resulting profile (depth, peak surface concentration, and Rsheet) can be tailored through laser pulse energy density. This technology adds only one step and is commercialized by Manz. In addition, centrotherm is working on a similar approach as well (Friess et al., 2010).

5.2.6 Laser doping via laser chemical processing and NiAg light-induced plating One of the selective emitter approaches developed at Fraunhofer ISE, now further studied at RENA, is based on simultaneous ablation of the PECVD SiNx:H layer and melting of the emitter layer underneath the ablated region (120 Ω/sq) using a liquid-guided laser beam (laser chemical processing) (Kray et al., 2010). The liquid contains P atoms serving as P-source and heavy doping is reached after recrystallization of the molten Si. The technology enables self-aligned light-induced plating of the front contact, e.g., via Ni and Ag. Only one extra step is added and plating allows for thinner, highly conductive grid lines compared to screen-printed contacts.

5.2.7 Laser doping and plating University of New South Wales, Australia, developed a process similar to the one described in the previous section starting with a 100–120 Ω/sq diffusion (Tjahjono et al., 2008). Instead of laser chemical processing, the doping source can be, e.g., phosphoric acid deposited on the wafer prior to laser State-of-the-Art Industrial Crystalline Silicon Solar Cells 49 doping. It allows for self-aligned plating of the front contacts as well. Both processes insert extra steps after firing of the Al-BSF, which can therefore be optimized independently of the front contact. In this approach, a plated Ni/ Cu/Ag front contact stack is used. Two extra steps are added, and the approach allows for thinner, highly conductive grid lines as well. Roth & Rau are working on commercialization.

5.2.8 Effect of encapsulation on blue response Selective emitters show an increase in j–V parameters on solar cell level as demonstrated above, but the enhanced performance has to occur on module level after encapsulation under glass as well. As part of the gain in short wavelength IQE might be lost due to absorption after encapsulation, experiments have been carried out to investigate this effect. In Fig. 1.24, transmission curves for pure module glass and EVA under module glass are shown. It can clearly be seen that the EVA starts to limit the transmission at around 380 nm, whereas the module glass transmits significantly shorter wavelengths. This negative effect of EVA concerning transmittance could be overcome by the use of alternative materials like silicones (Ketola et al., 2008; Ohl and Hahn, 2008). As shown in Fig. 1.24, the transmission of silicone under module glass is almost identical to the curve of the module glass alone. Therefore, encapsulation of solar cells under glass using silicones would provide a better use of the short wavelength photons.

100

40 Module glass

Transmission (%) EVA on module glass 20 Silicone on module glass

0 400 600 800 1000 1200 l (nm) Figure 1.24 Transmission of module glass, EVA under module glass, and silicone under module glass. Whereas EVA limits transmittance below 380 nm, silicone has a better transmission of short wavelength photons. Data from Ohl and Hahn (2008). 50 Giso Hahn and Sebastian Joos

1.0

0.8

0.6 IQE 0.4 Hom. emitter (EVA) Sel. emitter (EVA) 0.2 Hom. emitter (silicone) Sel. emitter (silicone)

0.0 300 350 400 450 500 550 600 l (nm) Figure 1.25 Effect of encapsulation of homogeneous and selective emitter solar cells under EVA and silicone. The shaded area visualizes the additional loss of selective emitter solar cells (<0.1 mA/cm2). Data from Hahn (2010).

Better transmission of silicone results in a calculated gain in jsc of almost 0.4 mA/cm2 as calculated in Ohl and Hahn (2008) for state-of-the-art homogeneous emitter cells. This effect is demonstrated in Fig. 1.25 where a homogeneous emitter cell and a cell with selective emitter using etch-back technology are compared after encapsulation under EVA and under silicone. The loss in jsc, calculated by the difference in IQE of the silicone and EVA curves under AM1.5g (around 0.4 mA/cm2 for the homogeneous emitter cell) is only slightly higher for the selective emitter cell. The difference between losses of the selective and homogeneous emitter cell can be roughly visualized by the shaded area and amounts to less than 0.1 mA/cm2. There- fore, it can be concluded that the additional loss of selective emitter cells by encapsulation under EVA compared to homogeneous emitter cells is almost negligible.

5.2.9 Efficiency potential of selective emitters Interestingly, the efficiency potential for all selective emitter technologies seems to be very similar. In 2010, all technologies presented above reached efficiencies around 19.0% on large area Cz-Si solar cells with full area Al-BSF (Hahn, 2010). Compared to standard state-of-the-art solar cells with homogeneous emitter in 2010, improvement in efficiency on cell level using a selective emitter was around 0.5%abs, with an increase in Voc of around 10 mV. This can be explained by the fact that for this cell concept the front State-of-the-Art Industrial Crystalline Silicon Solar Cells 51 side is not the limiting region anymore. The efficiency of these cells is mainly limited by the recombination at the rear side (full area Al-BSF). Since 2010 a lot of effort was put into development of new Ag pastes, allowing for higher Rsheet to be contacted. Therefore, part of the motivation for application of selective emitter structures is weakened. On the other hand, some benefits of selected selective emitter technologies, like broad firing windows and better electrical yield due to deeper emitters remain and are still of importance. In addition, the formation of selective emitters might be of higher interest again when the solar cell design on the rear side is changed toward local contacts, decreasing recombination at the rear and putting emphasis again toward the front side quality.

6. INDUSTRIAL PERC-TYPE SOLAR CELLS

In addition to the improvements of the standard process discussed in the last section, a lot of effort is spent to implement more radical changes in cell architecture into the industrial mass production cell process. After optimization of the emitter quality described in the last section, the full area Al-BSF is now the main reason for limiting efficiency of cells fabricated from high quality Si material to values below 20% in mass production. Therefore, different concepts based on locally contacted rear sides in combination with a dielectric passivation of the largest part of the rear side area have been developed. Some concepts available for this purpose have been introduced first on lab-type solar cell processes and are the passivated emitter and rear cell (PERC, Blakers et al., 1989), the passivated emitter and rear locally diffused (PERL, Wang et al., 1990), and the passivated emitter rear totally diffused (PERT, Zhao et al., 1999) concepts (Fig. 1.26). All of these concepts have in common that most of the rear side area is dielectrically passivated (by thermally grown SiO2 in the original lab-type processes) and only local contacts are applied (mainly by photolithography for lab-type processes). The differences between the mentioned concepts are • PERC: metal contact directly on Si with moderate base doping • PERL: locally diffused highly doped area (BSF) in the contacted areas only • PERT: rear side totally diffused (full area BSF) with local contacts. For p-type solar cell concepts, the BSF (locally for PERL or full area for PERT) was first applied via B diffusion. For all concepts, recombination of charge carriers at the rear is drastically reduced compared to the standard process described in the sections above. Efficiencies of up to 25% could be 52 Giso Hahn and Sebastian Joos

Ag SiNx:H

n n+ p-Si

Al2O3/SiNx:H Al Ag SiNx:H

n n+ p-Si p++

Al2O3/SiNx:H Al Ag SiNx:H

n n+ p-Si p++

Al2O3/SiNx:H Al

Figure 1.26 Schematics of Al2O3/SiNx:H passivated rear side concepts with selective emitter: PERC (top), PERL (center), and PERT (bottom). reached on small area cells (Zhao et al., 1999), demonstrating the efficiency potential. The drawback of these cell concepts was for a long time that the cell process was very complex, adding many extra steps and therefore calling its industrial realization into question. In recent years, a lot of effort was put into development of alternative processing steps allowing for an industrial realization of a passivated rear side solar cell concept. In the following, progress toward implementation of these cell concepts into industrial production will be reported.

6.1. Dielectric rear side passivation For lab-type solar cells, rear side passivation was originally realized by long high temperature thermal oxidation steps to grow SiO2 layers in the range of 100 nm thickness at temperatures above 1000 C. This type of processing is thought not to be applicable in industrial mass production for reasons of throughput and material issues.23 Therefore, alternative passivation schemes

23 Many c-Si materials cannot withstand temperatures above 950 C because of the underlying defect structure. This is especially important for mc-Si material. State-of-the-Art Industrial Crystalline Silicon Solar Cells 53 had to be developed for industrial solar cell processing. Surface passivation of p-type surfaces can either be reached by chemical passivation, or by field effect passivation with layers having negative fixed charges close to the dielectric/silicon interface. In this case, the negative fixed charges attract holes (majority carriers) and repel electrons (minority carriers), reducing the recombination probability. In principle, also positive fixed charges at the interface as present, e.g., in SiNx:H layers can reduce surface recombination for p-type silicon, provided that the charge is high enough to bring the semiconductor in the regime of inversion (then attracted electrons show a higher density close to the surface than holes, again decreasing recombination). But it could be shown that such a concept based on an inversion layer is not very well suited for solar cell production (Agostinelli et al., 2005; Finck von Finckenstein et al., 2000), as there is a high probability that the inversion layer will be shunted during metallization (Dauwe et al., 2002).

6.1.1 Al2O3 layers It is known for quite some time that Al2O3 layers can offer a good surface passivation effect for PV application (Hezel and Jaeger, 1989; Ja¨ger and Hezel, 1985). The reason is the existence of a high amount of fixed negative charges at the dielectric/silicon interface, allowing for excellent passivation of p-type silicon surfaces (Agostinelli et al., 2006). The field effect passivation of the Al2O3 by negative fixed charges has to be activated by temperatures of around 400 C or higher (Hoex et al., 2006). Possible ways for deposition of Al2O3 layers are, e.g., atomic layer deposition (ALD; Agostinelli et al., 2006; Hoex et al., 2006), PECVD (Miyajima et al., 2008), or APCVD (Black and McIntosh, 2012). First large area solar cells with Al2O3 rear side passivation have been reported in 2009 (Vermang et al., 2009), but the layers showed poor stability during the high temperature firing step needed for screen-printing contact formation.

6.1.2 Al2O3/SiNx:H stacks The refractive index of Al2O3 is in the range of 1.6–1.7 (Hoex et al., 2006) and therefore not optimal for optics. A possible way to improve optics is the application of a stacked layer system with a thin layer of Al2O3 and a thicker layer of PECVD SiNx:H with refractive index of around 2.0. This allows for better optics as well as a sufficiently good firing stability (Dingemans et al., 2009). Positive fixed charges of the SiNx:H layer have no negative effect on the passivation quality of the Al2O3, as long as the Al2O3 layer exceeds a minimum thickness of around 5 nm (Richter, 2014; Schmidt et al., 2010). 54 Giso Hahn and Sebastian Joos

First large area solar cells with Al2O3/SiNx:H rear side passivation have been published in 2010 (Lauermann et al., 2010), and reasonably good surface passivation after the firing step could be demonstrated.

6.1.3 SiO2/SiNx:H stacks An alternative to Al2O3/SiNx:H stack systems is the use of SiO2/SiNx:H stacks. Here, the SiO2 layer can, e.g., be grown by thermal oxidation or by PECVD, the latter allowing that both layers can be deposited in a single PECVD run. The SiO2 layer can shield the positive fixed charges of the SiNx:H and provides a good chemical passivation, subject to the condition that a sufficient amount of H is present at the SiO2/Si interface. Therefore, SiO2/SiNx:H stack systems usually gain in surface passivation after the firing step, whereas the passivation quality of Al2O3/SiNx:H stacks is reduced after 24 firing. Nevertheless, the absolute overall passivation quality of Al2O3/SiNx: H stacks after firing is usually still higher than for SiO2/SiNx:H stacks.

6.2. Formation of local rear contacts For local rear contact formation, the formed dielectric layer has to be opened locally. Different approaches have been developed for this purpose: • Local firing of the metal through the dielectric layer • Local application of an etching medium followed by local etching of the dielectric layer • Local ablation of the dielectric layer, e.g., by laser.

6.2.1 Laser fired contacts An elegant way for realizing a local rear side metallization scheme is the formation of laser fired contacts (LFC; Schneiderlochner€ et al., 2002). Here, the full area Al is deposited on the not yet opened dielectric layer. Afterward, a laser pulse is applied and locally melts the Al and removes the dielectric. Dur- ing cooling down, a contact between Al and c-Si is formed. This approach was first developed for evaporated or sputtered layers of Al, and it could be shown that during cooling down a thin Al-BSF is formed at least in the center of a LFC (Glunz et al., 2003). This approach could also be transferred to large area cells with screen-printed Al layers, with the LFC formation applied after the firing step (Hofmann et al., 2008).

24 Compared to layers where the passivation was activated by an annealing step at around 400 C. State-of-the-Art Industrial Crystalline Silicon Solar Cells 55

6.2.2 Etching paste A medium able to etch through the dielectric layer can be applied locally to the rear side of the wafer. This could, e.g., be realized via inkjet or screen- printing of a paste containing phosphoric acid as etching medium (Ba¨hr et al., 2007). Normally, a thermal treatment is needed after application of the etching paste to activate the chemical reaction, but the advantage is that the c-Si underneath the dielectric layer is not damaged.

6.2.3 Laser ablation Another approach is local ablation of the dielectric layer by using a short laser pulse. Here, the aim is to remove the dielectric layer locally, with minimal damage of the c-Si underneath. As the photon energy of the laser pulse should be absorbed mainly in the dielectric layer, picosecond (ps) laser pulses with wavelengths of 532 and 355 nm show very good results on planar surfaces (Engelhart et al., 2007; Hermann et al., 2010). High photon energy and density result in most of the pulse energy being absorbed in the dielectric layer, heating it up and evaporating it immediately. The short pulse length assures that the thermal heat wave does not penetrate deep into the c-Si underneath, restricting damage due to, e.g., thermal stresses to surface near regions. After local opening via etching or laser ablation, the Al paste is screen- printed on the rear. During firing in a belt furnace, the Al paste without glass frit does not etch through the dielectric, but forms only a local contact in the areas previously opened.

6.2.4 Contact patterns and void formation While the LFC approach uses points as contact geometry, for the laser ablation approach a line-shaped contact pattern is advantageous (Lauermann et al., 2011). This can be understood considering the different formation mechanisms of the local contacts. For LFC, during the very short laser pulse no significant diffusion of Si into the Al layer can occur. This situation changes when the contact is formed during the firing step in a belt furnace. Here, the situation is similar to the standard full area BSF formation. Above the eutectic temperature of Al/Si, Si is dissolved into the molten Al. As this dissolution can only occur locally in the opened areas, the amount of Si dissolved into the Al depends on the geometry of the opening (Urrejola et al., 2011). Depending on the temperature–time profile during firing, different amounts of Si can diffuse into the Al layer, saturating the molten Al in some areas, while others further away from the opening show lower Si 56 Giso Hahn and Sebastian Joos concentrations below the values predicted by the Al/Si phase diagram (Urrejola et al., 2010). During cooling down after firing, the BSF starts to form when Si is repelled from the Al/Si melt, again according to the phase diagram. For doing so, Si has to diffuse back to the local openings in the dielectric layer. Depending on the amount of Si dissolved, BSF formation might start at lower temperatures compared to the case of a full area BSF. For a detailed model of local BSF formation using laser ablation see Lauermann et al. (2013). To reduce the effect of Si being dissolved into Al only to values below the values according to the phase diagram at peak temperature, Si can be added to the Al paste, reducing this unwanted effect, and improving BSF formation (Rauer et al., 2011). It is often observed that after application of local Al rear contacts voids in the c-Si are formed in the opened areas (Meemongkolkiat et al., 2006). This observation is often explained by the Kirkendall effect (Kirkendall, 1942). These voids can be visualized by cross section microscopy or in a nondestruc- tive way, e.g., by scanning acoustic microscopy (Dressler et al., 2012). Not all voids have a direct negative influence on cell parameters. If a BSF is formed and is in contact with the Al layer, the void does no harm, whereas on the other hand a missing BSF is responsible for increased surface recombination, and a noncontacted BSF causes increased Rs. See also Riegel et al. (2012) for a review of screen-printing metallization on p-type Si.

6.2.5 Interconnection issues The designs for local Al contact formation described above all result in reduced recombination losses at the rear side. But for module interconnection ribbons have to be soldered to both rear and front contacts. On the rear there are no Ag pads as in the case of the standard full Al-BSF cells, and as Al paste is not solderable alternatives have to be found. One of them is the deposition of Sn stripes on the rear using ultrasonic soldering (also called tin pad; von Campe et al., 2012). On these pads interconnection ribbons can be soldered. Finally, it has to be mentioned that the cell structure with passivated rear side and locally alloyed Al contact is often referred to as PERC or industrial PERC structure. This might be somewhat misleading, as in the original publications of the PERC solar cell structure, the metal is directly contacting the base doping, without a highly doped p+ layer in between (Blakers et al., 1989). The structures described above would therefore more precisely fall in State-of-the-Art Industrial Crystalline Silicon Solar Cells 57 the category of PERL cells, with the difference that the p+ areas below the local metal contacts are not formed by (B) diffusion, but by (Al/B) alloying. Nevertheless, the term PERC structure is used more frequently.

6.3. Boron–oxygen related degradation With the front side improved using an optimized homogeneous emitter or a selective emitter structure and recombination at the rear side reduced by applying a passivation layer and only local contacts, τb sets the limit of efficiency. For mc-Si, the underlying defect structure (dislocations, grain boundaries, impurities, like, e.g., transition metals) can be identified as source for τb limitation. For higher quality Cz-Si, the formation of boron–oxygen (B–O) related defects plays a crucial role and is known for a long time (Fischer and Pschunder, 1973). These defects can form at room temperature in the presence of excess charge carriers, reducing τb significantly (Bothe et al., 2003; Weizer et al., 1979). When the solar cell is illuminated during operation, B–O related defects form within the first hours of operation of a solar module. The amount of defects formed scales linearly with substitutional boron concentration Bs (Glunz et al., 2001; Schmidt and Cuevas, 1999), and quadratically with interstitial O concentration Oi (Rein et al., 2003; Schmidt and Bothe, 2004). A fundamental lifetime limit in dependence of Bs and Oi could be found for the degraded state after formation of the B–O related defects (Bothe et al., 2005), restricting τb to

45 0:824 1:748 τb ¼ 7:675 10 Bs Oi : (1.39) The recombination active complexes become instable at temperatures exceeding 140 C in an anneal step in the dark, therefore the defect is called metastable (Rein et al., 2001). The behavior of the metastable B–O related defect can be understood by a conversion of two possible states into each other (annealed state and degraded state). In the annealed state, the complex shows negligible recombination activity, but in the presence of charge carriers it is transformed into the degraded state again, with a thermally activated reaction rate.

6.3.1 Strategies to deal with B–O related degradation For most standard solar cell concepts, bulk resistivities of around 0.5–1 Ω cm allow for the highest conversion efficiency. With Oi in standard Cz material 18 3 16 3 in the range of 10 per cm and Bs of 10 per cm , bulk lifetime calculated according to Eq. (1.39) results in only 9 μs for the degraded state. Therefore, 58 Giso Hahn and Sebastian Joos strategies to prevent this severe degradation effect have to be considered. One possibility is to avoid B as a dopant and the usage of another p-type dopant like Ga (Glunz et al., 2001). Unfortunately, Ga has a segregation coefficient much smaller than B, resulting in inhomogeneous doping and therefore variations in bulk resistivity with ingot height (Trumbore, 1960). Another possibility is to reduce formation of Oi during crystallization. This can, e.g., be achieved by applying an external magnetic field (Mosel et al., 2012). Very good results have been achieved (Glunz et al., 2001), but extra costs have to be considered. The easiest way is to use a lower Bs, as B is deliberately added to the melt. This reduces the efficiency potential in the annealed state, but allows for higher efficiencies in the degraded state after illumination. This approach is followed in industry today, as 2–5 Ω cm B-doped Cz material is normally used.

6.3.2 Regeneration of B–O related defects In 2006, it was discovered that the degraded state of the B–O related defect can be transformed into another state, which is stable under illumination and shows negligible recombination activity (Herguth et al., 2006a,b). This regenerated state can be reached by slightly elevated temperatures and the presence of excess charge carriers25 (Herguth et al., 2008). This new state of the B–O related defect is stable under illumination and room temperature, and therefore allows for high and stable efficiencies at operating conditions of a solar cell under illumination. The three states of the defect (annealed, degraded, and regenerated) are depicted in Fig. 1.27. For higher temperatures >200 C, the regenerated state can become unstable, and the B–O related defect is in the annealed state, from where it can degrade again in the presence of excess charge carriers at lower temperatures. It is still unclear if the reaction path is directly from state C to state A (destabilization in Fig. 1.27) or via state B. A model based on reaction constants has been pro- posed to explain and predict the observed kinetics of the defect in detail (Herguth and Hahn, 2010). It could be shown that the presence of H might have an influence for the regeneration reaction to occur ( Mu¨nzer, 2009). This could be verified, as test samples with unfired SiNx:H layers did not show a regeneration effect, while fired samples did (Wilking et al., 2013a). Also layers inserted between c-Si and SiNx:H acting as a barrier for H slow down the regeneration

25 Excess charge carriers can be generated via illumination or biasing of the solar cell. State-of-the-Art Industrial Crystalline Silicon Solar Cells 59

State A “annealed” (inactive) Destabilization

State C Degradation Anneal “regenerated” (inactive)

Redegradation

State B Regeneration “degraded” (active) Figure 1.27 Three state model of the B–O related defect. States A and C are recombination inactive, but only state C is stable under illumination (after Herguth et al., 2006b).

process, indicating that H in sufficient concentration in the c-Si bulk is a prerequisite for the regeneration process to work. The firing process plays a crucial role in adjusting the concentration of H in the final solar cell, strongly influencing the regeneration kinetics (Wilking et al., 2013b). Choosing the right firing parameters and regeneration conditions, the regeneration process can be carried out within seconds (Wilking et al., 2014), opening a path toward stable efficiencies for p-type B-doped material with high Oi and bulk resistivities in the range of 1 Ω cm.

6.4. State-of-the-art industrial PERC solar cells Efficiencies of 20.2% have been demonstrated on large area p-type boron doped Cz-Si in 2011 using both the LFC technology (Engelhart et al., 2011) as well as laser ablation for local opening of the dielectric and use of an etched-back selective emitter (Gassenbauer et al., 2013). Optimiza- tions led to efficiencies of up to 21.0% for the laser ablation approach in 2013 (Metz et al., 2014; Ramspeck et al., 2012). Applying the regeneration process, these high efficiencies are stable under illumination. The efficiencies reported here have been reached with screen-printed metallization on the front and rear side. Even higher efficiencies of up to 21.3% are possible when plating approaches are used (Metz et al., 2014). For some of these cells described above, a thin thermal oxide layer is placed 60 Giso Hahn and Sebastian Joos

between the emitter and the SiNx:H ARC, improving surface passivation. In addition, optimized emitter structures including an etch-back process are used (Metz et al., 2014). Interestingly, the described industrial PERC processes might be easier to apply to mc-Si instead of Cz-Si. This seems to be counterintuitive, as Cz-Si might contain fewer defects than mc-Si, resulting in higher material quality, and therefore the efficiency potential of Cz-Si should be more sensitive for a lower rear surface recombination. But without application of the regeneration process to deactivate the B–O related defects, Cz-Si of higher base resistivity is used to avoid strong degradation under illumination (see above). Therefore, the p+/p interface formed between the full area Al-BSF and the c-Si forms a steeper high/low junction in the case of high bulk resistivity, lowering seff for Cz-Si to values below the ones for 1 Ω cm mc-Si (see Eq. 1.36). This means that it is more difficult for Cz-Si of higher bulk resistivity to achieve lower rear surface recombination by dielectric passivation as compared to full area Al-BSF processing than for mc-Si of lower bulk resistivity.

7. SUMMARY AND OUTLOOK

Since 2000 the yearly shipment of PV cells and modules increased by a factor of 100, indicating the tremendous progress that could be achieved. Interestingly though, the standard industrial solar cell fabrication process did not change too much over the past one to two decades. B-doped p-type wafer based Si is still the material of choice, and screen-printing of metal pastes is used for metallization. Most of the processing steps have been optimized. Examples are the development of high Rsheet emitters, resulting in lower j0e. This was only possible by the parallel development of new Ag pastes, allowing the contacting of these emitter structures without increased contact resistivity. Advances have also been achieved in front side screen-printing, with typical Ag finger width in the range of 60–80 μm. Other examples are optimized front surface texture solutions and Al rear side pastes, co-doped with B. But also more disruptive changes have been implemented into the standard cell fabrication process. After introduction of PECVD SiNx:H as ARC and surface passivation layer, solar cell efficiencies could be significantly increased, especially for defect-rich mc-Si solar cells. In addition, with the introduction of acidic surface textures mc-Si now shows a significantly higher market share as compared to Cz-Si, due to lower wafer fabrication State-of-the-Art Industrial Crystalline Silicon Solar Cells 61 costs. Selective emitter structures have been introduced into mass production, especially when they fit into the layout of the cell process.26 Altogether, an evolutionary process development was chosen based on the very robust industrial standard process to minimize the risk of failure. Especially in times of very dynamic market growth this approach is favored, as existing solar cell lines can be duplicated adding only slight changes in equipment or solar cell design. For the introduction of passivated rear sides with local rear contacts (PERC-type structures), the cell process has to be adapted more drastically, as several processing steps have to be added or changed. It remains to be seen if this technology will be the next step to be introduced into the standard process, or if it will be a parallel approach. Last but not least, there are many other cell designs under development at the moment, all introducing more disruptive changes compared to the standard process. Some of them are already in industrial mass production (or close to), but most of them are currently used by only a handful of companies. Therefore, these technologies have not been part of this chapter. Examples are • n-type Si wafers • Rear contacted solar cells (van Kerschaver and Beaucarne, 2006) • Heterojunction emitter technology (Ballif et al., 2014). The reason why these very interesting technologies are under investigation is the fact that they currently show higher efficiency potential. Whether this can also result in a more cost-effective PV electricity generation compared to the standard process route still remains to be seen, as many other factors apart from efficiency play a crucial role (yield, stability of the process, location of PV installation). It is hard to predict which process route and cell design will be the most cost-effective in the future. But it can be assumed that within the next decades there will be several c-Si technologies followed in industrial production in parallel, with many of them being commercially successful in the mass market or for niche applications. Today solar cells and modules are mainly optimized according to STC (25 C, 1000 W/m2) as this measurement is fast enough to be included in a production line and states a good measure for high irradiance environ- ments. The field conditions for solar cells, however, may be quite different. Here, higher module temperatures, lower irradiations, and stray light

26 An example is the etch-back selective emitter process, if edge isolation is carried out wet chemically anyway. 62 Giso Hahn and Sebastian Joos conditions are often found. Therefore, an additional optimization of PV solar systems in terms of generated electricity under these conditions is desirable.

ACKNOWLEDGMENTS We like to thank H. Haverkamp and A. Zuschlag for proof-reading and the German government for continuous funding of projects over the past decades.

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Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells

Christophe Ballif*,1, Stefaan De Wolf*, Antoine Descoeudres*, Zachary C. Holman† *Photovoltaics and Thin-Film Electronics Laboratory, Institute of Microengineering (IMT), Ecole Polytechnique Fe´de´rale de Lausanne (EPFL), Neuch^atel, Switzerland †School of Electrical, Computer, and Energy Engineering, Arizona State University, Tempe, Arizona, USA 1Corresponding author: e-mail address: [email protected]

Contents

1. Introduction 74 2. Passivating c-Si Surfaces with a-Si:H 76 2.1 Recombination at surfaces 78 2.2 Physics of passivation 79 2.3 Deposition of high-quality a-Si:H films 79 2.4 Surface passivation on n- and p-type wafers 82 3. From Passivated Wafers to Complete Solar Cells 83 3.1 Wafer cleaning and texturing 83 3.2 Electron and hole collectors: Doped a-Si:H layers 85 3.3 Transparent conductive oxide layers 86 3.4 Metallization 90 3.5 Record cells 92 4. Losses in Silicon Heterojunction Solar Cells 95

4.1 Voc losses 95 4.2 FF losses 95

4.3 Jsc losses 96 5. Industrialization and Commercialization 99 5.1 General status 99 5.2 Material requirements 100 5.3 Temperature coefficient and energy yield 102 5.4 Metallization 103 5.5 Tools and production technologies 103 5.6 Production costs 107 6. Future Directions and Outlook 108 Acknowledgments 110 References 110

1. INTRODUCTION

Silicon wafer-based solar cells have dominated the photovoltaics market for decades and may well continue to do so for years to come. Several key factors explain the success of this technology: Silicon is a well-studied semiconductor with known optoelectronic properties; it is abundant and nontoxic, and the price of multicrystalline silicon has witnessed an unprec- edented drop in the last few years, partially because of a temporary production overcapacity, especially in Asia; and silicon solar cell technology has greatly benefited from the accumulated knowledge in semiconductor processing developed by the microelectronics community. An important strength of the current industrial silicon solar cell technology is its fabrication simplicity. Only a few steps suffice to fabricate a full device, where each step often fulfills several roles. Examples of this are the emitter diffusion process, which simultaneously getters impurities from the bulk of the wafer, and the metal contact firing through the silicon nitride anti-reflection coating, during which bulk hydrogenation of the wafer also occurs. A drawback of this simplicity is that further improvements in device performance must rely on the increasing sophistication of existing processes, while fundamental short- comings of the technology are hard to overcome. One such critical limitation is carrier recombination at the electrical contacts. Carrier recombination in silicon is a well-understood phenomenon and its minimization is a key factor in obtaining high-efficiency solar cells. We make a distinction between intrinsic recombination (Auger and band-to-band radiative recombination) and deep-defect-mediated recombination (Richter et al., 2013). Importantly, the latter type of recombination can in theory be completely eliminated by using perfect crystals, combined with an “ideal” solar cell architecture. This structure should feature perfectly passivated surfaces and contacts. Taking a 100-μm-thick wafer, such a solar cell would yield an open-circuit voltage (Voc) of about 770 mV (Richter et al., 2013; Tiedje et al., 1984). Note that, with perfect contacts, the Voc represents the energetic distance between the quasi-Fermi levels, which themselves express the density of excess charge carriers present in the material as a consequence of shining light on it. An important reason why the Voc can never equal the bandgap of the absorber—1.12 eV for crystalline silicon (c-Si) at room temperature—is not only the operating temperature of the device but also the aforementioned intrinsic recombination losses. Despite this, it is possible to come close to the 770 mV limit in real devices with excellent surface and contact passivation. Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 75

Defect recombination in the bulk of c-Si has been successfully combated in recent decades, as evidenced by the ever-increasing quality of silicon wafers on the market. Surface passivation has also improved: A number of dielectric passivation layers are available that can passivate p-type and n-type surfaces very well. These include materials like silicon oxides (Benick et al., 2011; Deal and Grove, 1965; Green, 2009; Schultz et al., 2004; Zhao et al., 1998), amorphous silicon nitrides (Lanford and Rand, 1978; Lauinger et al., 1996) and aluminum oxide (Agostinelli et al., 2006; Hezel and Jaeger, 1989; Hoex et al., 2006). Surface passivation can be accomplished in two fundamentally different ways: Either the surface defect states are removed, or the excess charge carriers are screened from the surface defects by an internal electrical field. The former is known as chemical surface passivation and can be obtained by, e.g., hydrogenation of these defects. The latter is known as field-effect passivation and is usually obtained by deposition of a fixed-charge dielectric on the surface under study, thereby repelling minority carriers inside the wafer from the defective surfaces. Positive-fixed-charge dielectrics repel the positively charged holes inside the semiconductor from the surfaces, and are ideally suited to passivate n-type surfaces. A prime example is silicon nitride, which has been used for the passivation of phosphorus-doped emitters in homojunction technology (Lanford and Rand, 1978; Lauinger et al., 1996). Negative-fixed-charge dielectrics repel negatively charged electrons from the surfaces and are used to passivate p-type surfaces. Here, the most studied dielectric is aluminum oxide, which is a material that received significant attention in the last few years as a potential passivating layer for the rear surface of homojunction solar cells (Agostinelli et al., 2006; Hezel and Jaeger, 1989; Hoex et al., 2006). Silicon nitride layers can be relatively easily integrated into existing c-Si solar cell processing, whereas the successful integration of aluminum oxide layers into industrial solar cells has proven to be more of a challenge. In all cases, contacts are needed to extract carriers from the solar cell. In standard homojunction solar cell technology, where the junction is fabricated by thermal diffusion or ion implantation, these contacts are usually defined by locally opening the dielectric passivating layers and making a direct Ohmic contact between the metal and semiconductor. Whereas the contact resistances of such contacts can be made low, the minority- carrier recombination occurring at their surfaces is of significant concern. This issue is fundamentally resolved by silicon heterojunction technology, where a thin, wider-bandgap layer is inserted between the metal contact and the optically active absorber (i.e., the silicon wafer). Qualitatively, this 76 Christophe Ballif et al. type of contact can be considered as a semi-permeable membrane for carrier extraction. On the one hand, it should prevent generated carriers from being collected instantaneously, as this will lower the energetic splitting of the quasi-Fermi levels and thus reduce the voltage of the device. On the other hand, the contacts should be sufficiently electronically transparent to guarantee that carriers can be collected at the device terminals before they recombine in the wafer due to intrinsic recombination processes. In principle, such contacts can be fabricated in several ways. Irrespective of the materials used, passivated contacts should feature excellent (chemical) surface passivation while also giving charge carriers an incentive to be driven toward either the electron- or the hole-collecting layers. In this chapter, the focus will be on heterojunction solar cells with layers fabricated from thin films of amorphous silicon or related materials.

2. PASSIVATING c-Si SURFACES WITH a-Si:H

For silicon wafer-based devices, thin films of hydrogenated amorphous silicon (a-Si:H) are particularly appealing candidates for passivated-contact formation. First, a-Si:H passivates c-Si surfaces very well, with electrical properties that are on par with the best dielectrics available. The passivation is mostly chemical, principally due to hydrogenation of surface states. Second, such layers can be doped relatively well, either n- or p-type, by adding the appropriate process gasses during deposition. This property enables the fabrication of contacts that are selective for either electron collection (when n-type a-Si:H is used), or hole collection (when p-type a-Si:H is used). This is of significant utility, as it allows us to not simply make passivating contacts but also to escape the need for a homojunction in the wafer. As the lateral conductivity of the a-Si:H layers is quite low, transparent electrodes that serve electrical and optical roles are usually applied. Another important reason for the success of such layers is the available knowledge regarding thin-film deposition technology. Whereas silicon homojunction solar cell technology has benefited greatly from developments taking place within microelectronics, silicon heterojunction technology benefits from the flat-panel and thin-film solar cell industries, which have developed planar deposition technology with remarkable uniformity over very large surfaces and with high throughput. Depositions over areas of several square meters coupled with nanometric precision are commonplace. In 1974, Fuhs et al. were the first to study the interface between a c-Si wafer and an a-Si:H film (Fuhs et al., 1974). A few years later, Pankove Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 77 reported on the surface passivating properties of such thin films (Pankove and Tarng, 1979). In the early 1980s, a new type of tandem silicon solar cell was reported by Hamakawa et al., consisting of a multicrystalline silicon bottom cell and an a-Si:H top cell, the so-called Honeymoon cell. As the emitter of the bottom cell was made from a-Si:H as well, this is likely the first solar cell incorporating a silicon heterojunction for the emitter formation (Hamakawa et al., 1983; Okuda et al., 1983). Of notable interest is that a-Si:H layers are usually deposited at temperatures below 200 C. This makes the heterojunction concept also particularly attractive to fabricate emitters on substrates that would not withstand the temperatures usually involved in homojunction solar cell fabrication. This point gave Panasonic (Sanyo at the time) the motivation to incorporate silicon heterojunctions into their thin-film multicrystalline silicon solar cells in the late 1980s (Taguchi et al., 1990). To study the properties of this a-Si:H/multicrystalline silicon junction, devices were also fabricated on regular silicon wafers. The first devices resulted in efficiencies close to 12% and were fabricated by direct deposition of a thin boron-doped a-Si:H(p) emitter on an n-type c-Si wafer. These cells were identified to suffer from substantial interface recombination. A first significant advance was made when a thin intrinsic a-Si:H layer was inserted between the doped emitter and wafer to alleviate this issue. This is the so-called heterojunction with intrinsic thin-layer (HIT) structure, which increased the efficiency up to 14.5% (Taguchi et al., 1990). A second step toward high-efficiency solar cells was made by using a similar heterostructure as a passivating back (electron-collecting) contact, bringing the efficiency to over 18% (Tanaka et al., 1992). From this, it was clear that both electron- and hole-collecting contacts need to be passivated, and both can be fabricated by planar heterojunction technology. In Fig. 2.1, we sketch an a-Si:H/c-Si heterojunction solar cell with front and rear intrinsic a-Si:H buffer layers, as developed by Panasonic, and its band diagram. In the last 20 years, Panasonic has increasingly refined the heterojunction concept, even though the essential characteristics remained unchanged. Specific attention has been paid to improve surface passivation, lower the optical losses, and increase the fill factor (FF) of the devices. This sustained effort found its culmination in a reported solar cell efficiency of 24.7% in 2013 for a large-area (>100 cm2) device (Taguchi et al., 2014), using a silicon wafer less than 100 μm thick. This device featured a Voc of 750 mV, a value approaching the theoretical limit, underlining the particular appeal of this technology. With the interdigitated-back-contact configuration, the same company reduced the current losses at the front of the cell 78 Christophe Ballif et al.

Ag TCO a-Si:H (p+) E a-Si:H (i) − c-Si (n) h+ e

a-Si:H (i) a-Si:H (n+) TCO x Ag

Figure 2.1 Schematic of a silicon heterojunction solar cell on an n-type wafer, including its band diagram. Adapted with permission from De Wolf et al. (2012a). while maintaining high Voc and FF. Panasonic recently reached an efficiency of 25.6% for a 144 cm2 device, the highest 1-sun efficiency ever reported for a silicon solar cell (Panasonic, 2014a).

2.1. Recombination at surfaces In a working solar cell, generated carriers are collected at the relevant contacts or they recombine. Whereas the former process constitutes the external device current, the latter is purely a loss mechanism. Prior to either of these two processes occurring, the generated carriers reside in the material, where they contribute to the voltage of the device. In open-circuit conditions, obviously no current flows, and thus the Voc is directly linked to carrier recombination processes. Microscopically, in c-Si, bulk recombination is usually caused by deep defects, which originate from impurities or crystal defects. Such recombination is usually described by Shockley–Read–Hall recombination statistics. Deep-defect recombination is also of serious concern at clean silicon surfaces. The reason for this is the fact that the silicon lattice consists of covalent Si–Si bonds that must be broken at the surface. Depending on the precise surface orientation, each silicon surface atom will feature either one (as on the (111) silicon surface) or two (as on the (100) silicon surface) dangling bonds. These clean surfaces are often not stable and may reconstruct into lower-energy configurations, which could feature, e.g., Si–Si dimer bonds. The remaining silicon dangling bonds need to be passivated, however. The silicon dangling bond is a so-called amphoteric defect, which implies that it can have three different charge states. In its neutral state, it contains a single electron and its energy level is approximately at midgap, halfway between the valence band maximum and conduction band Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 79 minimum. In this state, it can give up its electron (positively charged state) or it can host a second electron (negatively charged state). The ease with which the dangling bond can move between these states, accepting either electrons from the conduction band or holes from the valence band, explains its high recombination activity and the need for surface passivation.

2.2. Physics of passivation The microscopic passivation mechanism of c-Si by a-Si:H is most likely closely linked to hydrogenation of surface states, where the hydrogen is supplied from the passivating film. For good passivation, it is necessary that the interface between the two materials be atomically sharp, i.e., that no epitaxial growth has takenplace (DeWolf andKondo,2007a;Olibetetal.,2010).Forsuchfilms, low-temperature annealing can also improve the passivation properties. For isothermal annealing, it was found that, irrespective of the precise deposition conditions, the electronic properties always obey stretched-exponential laws over the annealing time (De Wolf et al., 2008). Based on this, it could be argued that the passivation is to a significant extent driven by microscopic rearrangement of hydrogen close to the interface. Comparison with the bulk properties of a-Si:H indicated that the dominant defect responsible for recombination must be the same both in the a-Si:H bulk and at a wafer surface, and that this defect is likely the silicon dangling bond (De Wolf et al., 2012b).Quitegenerally,thepassivationpropertiesofa-Si:Hfilms mimic what occurs in their bulk, which includes effects such as light-induced degradation (De Wolf et al., 2011) and doping-induced defect generation (De Wolf and Beaucarne, 2006; De Wolf and Kondo, 2007b, 2009).

2.3. Deposition of high-quality a-Si:H films Since the passivating intrinsic a-Si:H layers are of such importance for the final device, great care has to be taken during their deposition. In particular, damage of the c-Si surface during deposition has to be limited as much as possible. The most common way of depositing these very thin layers is by plasma-enhanced chemical vapor deposition (PECVD), using silane (SiH4)—often mixed with H2 for an additional source of hydrogen—as the gas precursor, in a capacitively coupled parallel-plate reactor configuration. 13.56 MHz radio frequency (RF) power is generally used for plasma excitation (De Wolf and Kondo, 2007a; Munoz et al., 2012; Schulze et al., 2010; Strahm et al., 2010a), but very high frequencies (VHF) of 40.68 MHz or even 70 MHz have also been applied with success for c-Si 80 Christophe Ballif et al. surface passivation (De Wolf et al., 2008; Descoeudres et al., 2013; Mueller et al., 2010). VHF PECVD is known to have two main advantages over the more common RF PECVD: higher deposition rates due to enhanced dis- sociation of silane molecules, and reduced ion bombardment on the substrate due to lower sheath voltages (Shah, 2010). The effect of ion bombardment during film growth on the passivation quality is still under debate, but it likely creates defects in the a-Si:H matrix, reducing the a-Si:H/c-Si interface passivation (Illiberi et al., 2011). Device-grade PECVD passivating a-Si:H films are typically deposited at temperatures of approximately 200 C, in the 10–100 Pa pressure range, and in the 10–100 mW/cm2 power density range. Even if thickness inhomogeneity issues can arise in large-area (>1m2) PECVD systems (even more critical when using VHF rather than RF) (Howling et al., 2004; Schmitt et al., 2002), a-Si:H depositions have already been demonstrated in industrial systems with excellent uniformity and passivation (Strahm et al., 2010b). The development of such industrial reactors dedicated specifically to silicon heterojunction solar cell technology undoubtedly benefited from the knowledge gained in recent years by the thin-film transistor and thin-film silicon solar cell industries (Shah et al., 2013). Other relatively similar deposition techniques have also been investigated for c-Si passivation with a-Si:H, such as direct-current PECVD (Das et al., 2008), inductively coupled PECVD (Zhou et al., 2012), hot-wire (or catalytic) CVD (Gielis et al., 2008; Schuttauf et al., 2011; Wang et al., 2010), electron-cyclotron-resonance CVD (Maydell et al., 2006), DC saddle-field glow discharge (Bahardoust et al., 2010), and expanding thermal plasma (Illiberi et al., 2010). Although some of these techniques can provide passivation results comparable to those obtained by RF PECVD (Schuttauf et al., 2011), the latter remains the most widespread technology at present, mostly because of its well-proven usability at the industrial scale. Although the passivation properties of an a-Si:H film can generally be improved by thermal annealing (De Wolf et al., 2008; Schulze et al., 2009), excellent passivation with layers in their as-deposited state appears nevertheless to be important, mostly for layer stability over time, but also for practical reasons in the case of an industrial fabrication process. It has been reported that device-grade, passivating intrinsic a-Si:H layers are deposited in plasma regimes close to the amorphous-to-(micro-)crystalline transition (Descoeudres et al., 2010); this material is similar to that required for the intrinsic layer in thin-film silicon solar cells (Collins et al., 2003; Guozhen et al., 2000). Such layers are generally hydrogen rich, and show a low bulk Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 81 defect density related to a hydrogen bonding configuration dominated by monohydrides (Si–H) (Kroll et al., 1996, 1998). In silane-based discharges, this transition is unequivocally determined by the actual SiH4 concentration in the plasma, cp ¼c(1D), with c the input SiH4 concentration and D the SiH4 depletion fraction induced by the discharge (Strahm et al., 2007). With the aid of plasma diagnostics such as infrared absorption spectroscopy or optical emission spectroscopy, it is indeed found that the best as-deposited a-Si:H passivating films are obtained with low-cp plasmas, which correspond precisely to amorphous regimes close to the crystalline transition (Strahm et al., 2007). These regimes can be produced either with SiH4 plasmas highly diluted with H2 (Gogolin et al., 2012) or with highly depleted pure SiH4 plasmas (Descoeudres et al., 2010; Howling et al., 2010). This result underlines again the clear role of atomic hydrogen, coming from H2 or SiH4, in the microscopic passivation mechanism at the a-Si:H/c-Si interface. As stated above, epitaxial growth has to be avoided in order to have high passivation quality. To work in regimes close to the amorphous-to-crystalline transition is therefore not without risk. A possible way to further approach the transition without epitaxial growth on the c-Si surface is the use of H2 plasma treatments, either during (Descoeudres et al., 2011)orafter(Mews et al., 2013) the a-Si:H deposition. Such treatments have several impacts on the deposited a-Si:H material. Depending on the plasma treatment conditions, one observes a modification of the material structure (either increased disorder in the silicon network (Descoeudres et al., 2011) or improved film quality, leading to better film stability with regard to light-induced degradation (Sakata et al., 1993)), and an increase in the hydrogen content (Fig. 2.2), which also widens the bandgap. Although treated a-Si:H films may be more disordered and can contain more recombinative defects than untreated films, the passivation quality of the crystalline substrate is generally greatly improved due to the increased hydrogen content in the film. Hydrogen diffuses from the bulk of the very thin a-Si:H layer toward the a-Si:H/c-Si interface, efficiently passivating silicon dangling bonds at the interface. Note that H2 plasma treatments before a-Si:H growth, i.e., directly on the bare c-Si surface, lead to the creation of persistent defects at the surface and thus to reduced passivation quality afterwards (Schuttauf et al., 2011). More generally, the impact of any plasma species impinging on the bare surface at plasma ignition is, to some extent, detrimental to surface passivation (Neitzert et al., 1993). This damage created at the very early stage of the plasma deposition process is then partly recovered by the passivating effect of the deposited film itself. Similarly, a prolonged H2 treatment on a 82 Christophe Ballif et al.

SiH

SiH 0.8

0.7 Standard layer 0.6 H2 treated layer

0.5

0.4

Absorbance 0.3

0.2 )

(O 0.1 2

SiH 0 2300 2200 2100 2000 1900 1800 Wavenumber (cm−1) Figure 2.2 Absorption spectra measured by attenuated total reflectance Fourier- transform infrared spectroscopy of 15-nm-thick H2-plasma-treated and untreated a-Si:H layers deposited on (111) c-Si wafers. Spectra are deconvoluted with two Gaussians, centered at 2000 cm1 (monohydride bonds, stretching mode) and at 2080 cm1 (higher hydrides bonds, stretching mode). Reproduced with permission from Descoeudres et al. (2011). deposited a-Si:H passivating film can lead to undesirable excessive film etching. Under a critical thickness, the etched passivating layer does not ade- quately protect the c-Si surface from the H2 plasma, and defects can also be created underneath at the a-Si:H/c-Si interface (Geissbu¨hler et al., 2013). H2 plasma treatments have also been demonstrated as a dry method to clean the wafer surface and thus a possible alternative to the traditional hydrofluoric acid (HF) dip before PECVD (Martı ´n et al., 2004).

2.4. Surface passivation on n- and p-type wafers Although p-type c-Si is the standard material for diffused-junction solar cells and therefore largely dominates current industrial photovoltaics production, n-type monocrystalline silicon appears to be the best candidate for high- efficiency solar cells and is foreseen to increase its share in c-Si photovoltaics production in the coming years (International Technology Roadmap for Photovoltaics, 2013). Indeed, the minority-carrier lifetime in the c-Si bulk is usually higher in n-type than in p-type material because most metallic Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 83 point defects have larger capture cross sections for electrons than for holes (Macdonald and Geerligs, 2004). Moreover, Czochralski (CZ) n-type wafers do not suffer from light-induced degradation, as is the case for CZ p-type wafers when a boron–oxygen or boron–iron complex is present (Lagowski et al., 1993; Schmidt and Cuevas, 1999). Therefore, considering only the basic bulk properties of c-Si wafers, n-type material is better suited to reach high conversion efficiencies. Regarding surface passivation with a-Si:H, fundamental differences exist between n- and p-type wafers. As with bulk defects, the capture cross sections of surface defects, i.e., dangling bonds, are larger for electrons than for holes (Olibet et al., 2010). The same phenomenon occurs for surface defects with thermally grown silicon dioxide (SiO2) passivation (Aberle et al., 1992), even though the asymmetry in the respective cross sections is much more pronounced in this case than for defects at the a-Si:H/c-Si interface. As a result, the injection-dependent minority-carrier lifetime curves are different for the n- and p-type cases (Fig. 2.3). A significant drop in lifetime is observed at low injection on p-type wafers because electrons (the minority carriers in p-type c-Si) are more easily lost at the interface than holes via defect-assisted recombination. This behavior cannot be attributed to bulk defects, because high-quality float-zone (FZ) wafers were used in this experiment. Notably, this drop in lifetime in the p-type case has a detrimental effect on the FF of completed solar cells (see Section 4.2). Aluminum oxide (Al2O3) layers are better adapted to p-type wafer passivation than a-Si:H. Higher low-injection lifetimes are obtained with such layers, due to the negative fixed charge present in this material (Hoex et al., 2008). At high injection, on the other hand, the lifetimes are similar in both cases, and are limited by unavoidable Auger and radiative recombination. Thus, there is the potential to reach very high Voc values, characteristic of the excellent surface passivation of silicon heterojunction solar cells, on both n- and p-type wafers, as shown by the relatively similar implied Voc values (Fig. 2.3). Such high-lifetime cells do indeed reach high injection at open circuit under 1-sun illumination.

3. FROM PASSIVATED WAFERS TO COMPLETE SOLAR CELLS 3.1. Wafer cleaning and texturing For silicon heterojunction solar cells, the wafer of choice is usually a monocrystalline silicon CZ wafer that is phosphorus doped and has (100) surface 84 Christophe Ballif et al.

Auger + radiative recombination limits

10−2 MPP

1 sun

V Implied oc = 730 mV V Implied oc = 732 mV

Lifetime (s) 10−3 MPP 1 sun

n-type p-type

1014 1015 1016 Minority carrier density (cm−3) Figure 2.3 Minority-carrier effective lifetimes of silicon heterojunction solar cell precursors (textured wafers passivated with co-deposited i-n and i-p a-Si:H stacks) on n- and p-type 4 Ω cm FZ wafers, measured with a quasi-steady-state photoconductance system (Sinton and Cuevas, 1996). The injection levels corresponding to 1-sun illumination are marked by solid arrows, and the corresponding implied Voc values are given. The injection levels corresponding to the maximum power points (MPPs) of the finished devices (under 1-sun illumination) are marked by the dashed arrows. Combined radiative and Auger recombination limits are shown by the solid lines (Richter et al., 2012; Schlangenotto et al., 1974). Reproduced with permission from Descoeudres et al. (2013). orientation. Its thickness may vary from about 200 to 100 μm(Taguchi et al., 2014), and it is usually pseudo-square with 600 sides if used for industrial silicon heterojunction solar cells (Papet et al., 2013). Monocrystalline silicon is usually preferred to multicrystalline silicon because of the higher carrier lifetimes usually associated with the monocrystalline material. In addition, the defined crystalline orientation of monocrystalline silicon allows for random-pyramid texturing. Such pyramids are formed by anisotropic etching of (100) surfaces, which reveals pyramids with (111) oriented facets (Bean, 1978). The bases of the pyramids are usually in the range of 5–10 μm, and the facets can feature steps but are usually quite flat. This flatness is of considerable importance to deposit thin films of equal thickness during subsequent processing. Such surfaces are almost impossible to obtain on multicrystalline wafers. One option for heterojunction technology on cast-silicon Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 85 wafers is to use quasi-monocrystalline silicon wafers ( Jay et al., 2012). How- ever, such wafers still require post-wafering processing steps such as gettering and hydrogenation to bring the bulk carrier lifetime to acceptable levels, casting doubt on whether this material is well suited for the fabrication of cost-effective silicon heterojunction solar cells. The preparation of the wafer surfaces prior to film deposition usually consists of several steps, some of which can be combined. First, damage caused by wafer sawing is removed in an alkaline solution. Next, the typical pyramidal texture is developed, also using an alkaline solution (Bean, 1978; Papet et al., 2006). This process is followed by wafer cleaning. Several methods exist for this purpose, but usually the philosophy is the same: The surfaces are wet-chemically oxidized, “encapsulating” impurities present on the surface, and then the oxide is stripped, usually in a dilute HF bath. Dur- ing this process, the impurities are removed, while the exposed silicon surface is terminated by hydrogen atoms. This yields chemical passivation, and can also stabilize the surface for some time following removal of the wafers from the chemical baths. Despite this, it is usually recommended to swiftly transfer the wafers into the deposition system.

3.2. Electron and hole collectors: Doped a-Si:H layers To give carriers generated in the silicon wafer an incentive to be collected, contacts specifically designed for the collection of electrons and holes must be designed. In principle, several approaches for forming such carrier- selective contacts exist; in silicon heterojunction solar cells, this is achieved by depositing thin doped a-Si:H layers on the passivation layers. For electron collection, a thin phosphorus-doped a-Si:H(n) film is used, whereas for hole collection a thin boron-doped a-Si:H(p) film is used. Though the doping efficiency for these materials may show some asymmetry— boron doping is well known to be difficult to achieve—the best devices reported to date rely on this type of contact. An integral part of these contacts is the transparent conductive oxide/metal stacks that are deposited on them. The precise contact formation, including the effect of the bulk and interface properties, has been the subject of intense study in recent years (Bivour et al., 2013, 2014a; Favre et al., 2013; Rossler et al., 2013). (Note that, apart from directly forming hole and electron collectors in silicon heterojunction solar cells, highly doped a-Si:H layers can also be used in c-Si solar cells as a phosphorus or boron dopant source for diffusion into the c-Si substrate (Seiffe et al., 2013)). 86 Christophe Ballif et al.

From a practical point of view, doped a-Si:H films are deposited like intrinsic films by PECVD with a SiH4 and H2 mixture (see Section 2.3), but with the addition of gases containing dopant precursors, highly diluted in H2. These are usually diborane (B2H6) or trimethylboron (B(CH3)3)forp-typea-Si:H,and phosphine (PH3) for n-type a-Si:H. As is well known in the thin-film silicon solar cell community, care has to be taken regarding cross-contamination if doped and intrinsic a-Si:H layers are successively deposited in a single PECVD chamber. Boron or phosphorus present in the a-Si:H on the reactor walls or substrate holder can be unintentionally incorporated into subsequent layers (Collins, 1988; Roca i Cabarrocas et al., 1989). Such contamination leads to poor p/i or n/i a-Si:H interface properties and is detrimental to the efficiency of thin-film silicon solar cells. Similarly, dopant contamination in the intrinsic a-Si:H passivation layers of silicon heterojunction solar cells leads to reduced passivation quality and thus Voc (more severely for boron than for phosphorus contamination). Dopants induce defects in the a-Si:H matrix, which act as recombination centers (De Wolf and Kondo, 2009; Korte and Schmidt, 2008). To circumvent these boron and phosphorus cross-contamination issues, several solutions have been developed (apart from using a multi-chamber PECVD system with chambers dedicated to i-, n-, and p-layer depositions): the deposition of a thick intrinsic coating on the reactor walls between layer depositions (Platz et al., 1997; Xiao-Dan et al., 2009), the use of a carbon dioxide (CO2)(Platz et al., 1997)orH2 plasma treatment (Cubero et al., 2011; Xiao-Dan et al., 2009), or the use of an ammonia (NH3)(Ballutaud et al., 2004) or water vapor flush (Kroll et al., 2004). Recently, more groups have started to investigate the use of specific electron- and hole-collecting materials that are not necessarily silicon based. For example, molybdenum oxide (MoOx) has been used in place of p-type a-Si:H as a hole-collecting layer. MoOx is a wide-bandgap material with a high work function. Therefore, it displays significantly higher transparency in the UV than p-type a-Si:H while maintaining the role of hole collector. With this approach, a short-circuit current density (Jsc) gain of about 1 mA/cm2 was obtained, compared to reference cells using p-type a-Si:H hole collectors, while maintaining a high Voc (Battaglia et al., 2014a,b).

3.3. Transparent conductive oxide layers In diffused-junction solar cells, the diffused emitter has a low sheet resistance of typically 50–100 Ω/sq because of its high doping density, mobility, and Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 87 thickness. In silicon heterojunction solar cells, the deposited emitter is made of low-mobility a-Si:H and it is only 5–10 nm thick. Consequently, low- resistance lateral transport is not possible in the emitter, and—as in thin-film solar cells that face the same obstacle—a transparent conductive oxide (TCO) layer is required at the front side to provide a low-resistance current path to the metal fingers. Its contribution to the lumped series resistance of the cell, and thus to power loss through Joule heating, is proportional to its sheet resistance, given by Rsheet ¼1/enμt, where e is the electronic charge, n is the free-electron density, μ is the electron mobility, and t is the layer thickness (Meier and Schroder, 1984). As in diffused-junction solar cells, sheet resistances of 50–100 Ω/sq are required for a typical finger spacing of approximately 2 mm if the TCO is not to contribute overwhelmingly to the lumped series resistance. The thickness of the layer is almost always fixed at approximately 75 nm since it then conveniently behaves as an excellent anti-reflection coating with a reflectance minimum at 600 nm. (This is possible because most TCOs have refractive indices of approximately 2, the geometric mean of air and silicon.) To reach the desired sheet resistance, instead of making the TCO layer thicker, the free-electron density is instead commonly tuned by adjusting the doping density. This approach is effective but has a negative side effect: Parasitic absorption of infrared light by free carriers increases with increasing free-electron density, reducing Jsc (discussed in detail later) (Holman et al., 2013a; Schroder et al., 1978). Conversely, increasing the electron mobility reduces sheet resistance and free-carrier absorption (Schroder et al., 1978). High-mobility TCOs are thus the holy grail of silicon heterojunction TCO research and are an important component of Panasonic’s success (Kinoshita et al., 2011; Taguchi et al., 2009, 2013), but obtaining such layers is challenging. For bifacial solar cells, the requirements for the rear TCO layer are similar to those for the front layer, but higher TCO sheet resistance can often be tolerated because the rear fingers are frequently closer together and the sheet resistance that is relevant to the lumped series resistance is that of the wafer and the rear TCO layer in parallel (assuming a front-emitter cell; for a rear- emitter cell, this is true for the front TCO layer). For a silicon heterojunction solar cell with full rear metallization, a TCO layer is not required for lateral transport—in fact, it is not clear that a TCO layer is required at all. While Bivour et al. demonstrated a 22.8% efficient rear-emitter cell without a rear TCO layer, in which a silver rear reflector was deposited directly on a highly doped p-layer (Bivour et al., 2012), a rear TCO layer is often included to reduce contact resistance (Holman et al., 2013a). In addition, Holman 88 Christophe Ballif et al. et al. showed that the rear TCO layer can play an important optical role if its thickness and carrier density are chosen correctly (Holman et al., 2013a,b, 2014). Near-bandgap p-polarized light that arrives at the rear TCO layer above the critical angle for internal reflection creates an evanescent wave that can be strongly absorbed in either the TCO layer itself or the subsequent metal reflector in the form of a surface plasmon polariton. By reducing the TCO free-electron density to n<1020 cm 3 and increasing the thickness to t>100 nm, both losses are suppressed, increasing the path length of weakly absorbed light in the wafer and thus increasing Jsc. For an interdigitated-back-contact (IBC) silicon heterojunction solar cell, the requirements for the rear TCO layers are the same as for cells with full rear metallization (though electrical contact needs to be made to both n- and p-type a-Si:H layers) and no TCO layer is required at the front. The most commonly used TCO in silicon heterojunction solar cells is indium tin oxide (ITO) with a tin oxide content of 5–20%. ITO is deposited by DC or RF sputtering in an argon atmosphere, and oxygen gas is added to tune the doping of the resulting layers through the density of oxygen vacancies (Buchanan et al., 1980; Holman et al., 2013a). Figure 2.4 shows the mobility and free-electron density, as determined from Hall-effect measurements, of ITO layers sputtered at room temperature using identical conditions but varying oxygen partial pressures. Two nominal film thicknesses were investigated (left column: 120 nm; right column: 290 nm) and the films were measured before and after curing at 200 C, which is the final step in silicon heterojunction fabrication after screen printing. In all cases, the mobility is relatively constant at a respectable but not impressive 20–40 cm2/Vs, whereas the free-electron density drops by roughly two orders of magnitude as oxygen vacancies are filled. To further increase mobility, researchers have explored doping indium oxide with tungsten (IWO) (Lu et al., 2013), hydrogen (IO:H) (Barraud et al., 2013; Koida et al., 2007, 2008, 2009, 2010, 2012), or both (IWOH) (Kobayashi et al., 2012). Lu et al. reported a mobility of 77.8 cm2/Vs for annealed IWO (Lu et al., 2013). Koida et al. achieved 120 cm2/Vs with IO:H sputtered in an atmosphere dosed with water vapor; they attributed the high mobility to an amorphous layer that crystallized to form larger grains upon annealing than can be achieved in as-deposited polycrystalline layers (Koida et al., 2007). This result was reproduced by Barraud et al., who fabricated a 22.1% efficient silicon heterojunction solar cell with an IO:H front TCO layer after they identified and solved a contact resistance problem between this layer and the screen-printed silver fingers (Barraud Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 89

Thin layers Thick layers 330 330 300 A B 300 270 270 240 240

120 120 Thickness (nm) 100 100

100 C D 100 80 80

/Vs) 60 60 2 40 40 30 30 20 20 Mobility (cm 10 10 1021 1021 EF ) − 3 1020 1020

1019 1019 Carrier density (cm

1018 1018 01234560123456

Plasma O2 content (%)

Uncured ITO 200 ЊC,10 min IO:H 200 ЊC,30 min Figure 2.4 (A) and (B) Thickness, (C) and (D) mobility, and (E) and (F) free-electron density of ITO and IO:H layers sputtered on glass before and after curing at 200 C. ITO films of the same nominal thickness were deposited with identical conditions but varying oxygen partial pressure. Characterization was performed with profilometry and Hall- effect measurements. et al., 2013). Figure 2.4 displays an IO:H layer in addition to ITO layers. Because of IO:H’s much higher mobility, switching from the best ITO layer in Fig. 2.4 to an IO:H layer with identical sheet resistance results in signif- 2 icantly less parasitic absorption (Fig. 2.5A) and a full 1.0 mA/cm gain in Jsc (Fig. 2.5B) (Barraud et al., 2013). Sputtered aluminum-doped zinc oxide (ZnO) is occasionally used in silicon heterojunction solar cells in place of indium-based TCOs (Maydell et al., 2006), but it is often too resistive (hard to achieve high free-electron densities) and does not have particularly high mobility. 90 Christophe Ballif et al.

15 A IO:H,43 Ω/sq ITO,38 Ω/sq 10

5 Absorbance (%)

0 100 B 80

40 IO:H, 40.0 mA/cm2 20 ITO, 39.0 mA/cm2

EQE and 1-reflectance (%) 0 400 600 800 1000 1200 Wavelength (nm) Figure 2.5 (A) Absorbance spectra of IO:H and ITO films on glass with the same thickness and sheet resistance. (B) External quantum efficiency (solid) and 1-reflectance (dashed) spectra of identical silicon heterojunction solar cells with IO:H and ITO front

TCO layers. Active-area Jsc values are given. Reproduced with permission from Barraud et al. (2013).

3.4. Metallization As in diffused-junction solar cells, the front fingers and busbars of silicon heterojunction solar cells are typically screen printed; a critical difference, however, is that the paste must be curable at temperatures below 250 Ctoavoid desorbing hydrogen from the a-Si:H layers, ruining the passivation. The two types of low-temperature paste are thermoplastic and thermoset (Zicarelli et al., 2010). The choice of paste depends on the required solderability for interconnection, on the targeted aspect ratio, and on the application (small cells for research or large-area devices). Compared to high-temperature fired paste (ρ¼3 μΩ cm), the best low-temperature paste typically reaches 8–10 μΩ cm. Silver nanopastes reach lower values but are prohibitively expensive and are hard to print thick enough. For large-area samples (>100 cm2), this factor of three in resistivity is a major drawback of silicon heterojunction technology. In other words, to reach similar finger-related losses as in diffused-junction c-Si cells, three times more paste has to be used. An easy approach to reduce paste consumption is to increase the number of Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 91 busbars. An increase from three to five busbars allows, to a first approximation, a decrease by 52/32 3 of the quantity of silver paste used for all fingers. A five- busbar silicon heterojunction cell can, hence, have finger-related FF losses similar to a three-busbar cell fired at high temperature. Three other approaches to reduce silver consumption are: 1. Low-temperature copper paste: Usually, such pastes include a low-melting- point alloy surrounding the copper particles. Some preliminary results were demonstrated (Tokuhisa et al., 2012), but cost and reliability data are not yet available. 2. Multi-wires: A promising approach is the use of multi-wire arrays, initially developed by Day4 Energy and now commercialized under the name “SmartWire.” This technology uses many wires in the place of busbars, and the wires are “soldered” to the finger paste during the lamination process (Fig. 2.6). The multi-wire approach is reported to reduce silver consumption to 40 mg per side of a 600 cell (Soderstrom et al., 2013), thanks to the relaxed requirements for finger conduction. The process is also compatible with bifacial cells. A 320 W module with a multi-wire approach was demonstrated recently (Kobayashi et al., 2013). Addition- ally, the multi-wire approach opens the possibility of other printing techniques for contacts, such as inkjet or offset printing, allowing a further reduction in silver usage (Hashimoto et al., 2013).

Figure 2.6 (A) Photograph of the multi-wire contacting scheme applied to the front fingers of a silicon heterojunction cell. (B) Scanning electron microscope image (prepared by using focussed ion-beam milling) of a nickel/copper finger. Shown below are energy- dispersive X-ray maps of copper, nickel, and indium, and a top-view scanning electron microscope image of the finger displayed above. Adapted with permission from Geissbühler et al. (2014). 92 Christophe Ballif et al.

3. Plated contacts: Several groups reported excellent results with copper- plated contacts (Geissbu¨hler et al., 2014; Hernandez et al., 2013; Munoz et al., 2012; Papet et al., 2013), including 24.2% efficient cells by Kaneka (Hernandez et al., 2013). The ITO layer that is usually used as the front contact and anti-reflection coating forms a good barrier against copper diffusion and offers a natural conductor for direct plating. In some cases, a nickel seed layer was reported to promote good adhesion of copper fingers on the solar cells (Geissbu¨hler et al., 2014), as illustrated in Fig. 2.6B. There are many approaches for creating the pattern or the seed layer, including inkjet printing of hot-melt wax (Hermans et al., 2013). In principle, copper metallization allows a large cost decrease compared to three- or even five-busbar silver-printed cells. However, it comes with the additional steps required to pattern the cell or to remove the mask.

3.5. Record cells In recent years, several academic groups have started to investigate silicon heterojunction solar cells. Meanwhile, companies have begun work on this technology as well. In the following tables we summarize the best reported results to date. Tables 2.1 and 2.2 show the best results on n- and p-type wafers, respectively. Table 2.1 clearly reinforces that this technology enables record values for Voc. Note as well that an increasing number of groups have 2 come very close to, or have even overcome, the 40 mA/cm mark for Jsc. Additionally, several groups have clearly demonstrated that FF>80% is possible, despite prior doubts. To put these results in perspective, in Table 2.3 we show the best diffused-junction solar cells. In Table 2.4, we show solar cells that belong to the family of heterojunction solar cells but have slight deviations such as a diffused front-surface field combined with a heterojunction rear emitter; SiOx tunnel oxides instead of intrinsic a-Si:H buffer layers; or epitaxial emitters combined with passivating heterostructure contacts. These results underline the fact that there are several approaches to fabricate high-efficiency silicon heterojunction solar cells. Table 2.5 shows the emerging trend of combining heterojunction technology with IBC solar cells. Because no shadow losses are present and no contacting structures are needed at the front, such a design may offer the ultimate solution that combines high Voc and high Jsc. In 2013, Sharp presented exciting results with an IBC structure, the precise processing of which remains undisclosed. These were followed in 2014 by the spectacular new Table 2.1 Device results on n-type c-Si wafers

η Voc Jsc FF Area Affiliation (%) (mV) (mA/cm2) (%) (cm2) Statusb Year Panasonic, Japan 24.7 750 39.5 83.2 101.8, IC 2013 (Taguchi et al., 2014) CZ (AIST) Kaneka, Japan 24.2 738.3 40.02 81.9 171.28, IC 2013 (Hernandez et al., 2013) CZ (ISE) RRR, Switzerlanda 23.14 736.7 38.64 81.3 239, – 2014 CZ EPFL and CSEM, 22.4 728 39.15 78.6 4, – 2014 Switzerland FZ (Geissbu ¨hler et al., 2014) CIC, Japan 22.3 733 37.28 81.8 243, – 2013 (Kobayashi et al., 2013) CZ AUO, Taiwan 22.26 724 37.5 81.97 238.9 – 2013 (Chen et al., 2013) CEA-INES, France 22.2 730 38.7 78.5 104 IC 2012 (Munoz et al., 2012) (ISE) Hyundai HI, Korea 21.1 721 36.6 79.9 220 – 2011 (Choi et al., 2011) SERIS, Singapore 21.1 702.2 38.2 78.6 1 – 2012 (Mueller et al., 2012) aThe cell was measured in a multi-wire configuration. bThe status column indicates whether the result was independently confirmed (IC) or appeared in a peer- reviewed publication (PR).

Table 2.2 Device results on p-type c-Si wafers

η Voc Jsc FF Area Affiliation (%) (mV) (mA/cm2) (%) (cm2) Statusb Year EPFL, Switzerland 21.4 722 38.4 77.1 4, IC 2012 (Descoeudres et al., 2013) FZ (ISE) NREL, USA 19.3 678 36.2 78.6 0.9, IC 2010 (Wang et al., 2010) FZ (NREL) Titech,a Japan 19.1 680 36.6 76.9 0.8, PR 2011 (Irikawa et al., 2011) FZ Ju¨lich,a Germany 18.5 664 35.7 78 0.76, PR 2013 (Ding et al., 2013) FZ aActive-area efficiency. bThe status column indicates whether the result was independently confirmed (IC) or appeared in a peer- reviewed publication (PR). Table 2.3 Best c-Si homojunction solar cells for p- and n-type c-Si wafers

η Voc Jsc FF Area Affiliation (%) (mV) (mA/cm2) (%) (cm2) Year UNSW, Australia 25 706 42.7 82.8 4, FZ, p 1998 (Zhao et al., 1996) SunPower, USA 24.2 721 40.5 82.9 155, CZ, n 2010 (Cousins et al., 2010)

Table 2.4 Notable exceptions on p- and n-type c-Si wafers

η Voc Jsc FF Area Affiliation (%) (mV) (mA/cm2) (%) (cm2) Type Statusa Year Fraunhofer 24.4 715 41.5 82.1 4, Diffused IC 2014 ISE, Germany FZ, front (ISE) (Heng et al., n emitter, 2013) SiOx tunnel oxide rear Fraunhofer 22.8 705 39.9 81.5 4, Diffused PR 2012 ISE, Germany FZ, front-surface (Bivour et al., n field, rear 2012) emitter

Silevo, USA 22.1 728 38.65 78.6 155, SiOx tunnel IC 2013 (Heng et al., CZ, oxide layers (Sandia) 2013) n IBM, USA 21.9 710 39.3 78.6 0.92, Epitaxial PR 2012 (Hekmatshoar FZ, emitter et al., 2012) p aThe status column indicates whether the result was independently confirmed (IC) or appeared in a peer- reviewed publication (PR).

Table 2.5 Device results for IBC silicon heterojunction designs

η Voc Jsc FF Area Affiliation (%) (mV) (mA/cm2) (%) (cm2) Year Panasonic, Japan 25.6 740 41.8 82.7 143, n 2014 (Panasonic, 2014a) Sharp, Japan 24.7 730 41.4 81.8 3.7, n 2013 (Koide et al., 2013) EPFL, Switzerland 21.5 724 39.9 74.5 9, FZ, n 2014 (Tomasi et al., 2014) LG, Korea 20.5 716 37.5 76.4 221, n 2013 (Lee et al., 2014) HZB, Germany 20.2 673 39.7 75.7 1, FZ, n 2011 (Mingirulli et al., 2011) CEA-INES, France 19.0 699 34.6 78.5 25, FZ, n 2012 (De Vecchi et al., 2012) Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 95 world record for a 1-sun c-Si cell from Panasonic, with 25.6% on 143 cm2. EPFL also recently demonstrated a high-efficiency IBC device based on simple processes (Tomasi et al., 2014). With additional efforts, improved FF values can likely be obtained, which would open interesting roads toward commercialization of this technology.

4. LOSSES IN SILICON HETEROJUNCTION SOLAR CELLS

4.1. Voc losses

Losses in Voc are due to recombination. Here, as already argued, the passivation of c-Si surfaces by a-Si:H films is quite remarkable, eliminating most of the surface states present. Intrinsic bulk recombination includes Auger recombination and radiative recombination. For the former, which strongly depends on carrier density, empirical expressions are available (Kerr and Cuevas, 2002) and recently have been revised (Richter et al., 2012). Miti- gation of the remaining defect recombination losses is one obvious way to obtain higher Voc values. Another is to use thinner wafers combined with excellent passivated surfaces. In this case, for the same AM1.5 1-sun illumination, the excess charge carrier density will be higher, thus yielding higher voltage. A concern when using thinner wafers, however, is infrared light management.

4.2. FF losses Causes of FF losses can be difficult to identify and suppress, since they come from different inter-dependent contributions. Shunts aside, FF losses in solar cells come from resistance to carrier transport (through each layer and across each interface) and from carrier recombination. Some of the transport-related contributions, like those from the a-Si:H/TCO interfaces, are specific to the structure of the silicon heterojunction solar cell. For the p-type a-Si:H/TCO interface especially, due to the n-type nature of most TCOs, the TCO work function and the doping level of the p-type a-Si:H layer critically affect transport via shifts in the band alignment (Bivour et al., 2013; Lachaume et al., 2013; Lee et al., 2013; Rossler et al., 2013). In particular, the p-type layer must be heavily doped and sufficiently thick so as not to be depleted by the adjacent TCO; thicker layers, however, increase blue parasitic absorption and reduce Jsc. This contact has to act as a band-to-band tunneling junction (Kanevce and Metzger, 2009), and is therefore also very sensitive to intra-band defect states in a-Si:H (Taguchi et al., 2008). Note that if the work functions of the metallization 96 Christophe Ballif et al. and the doped a-Si:H layer are well matched, the TCO layer can even be omitted without any loss in FF (Bivour et al., 2012). Another approach to relax the constraints linked to the aforementioned FF-Jsc tradeoff is to reverse the classical cell structure and to place the p-n junction at the rear of the cell (rear-emitter cell). This way, the TCO and p-type a-Si:H layers can be optimized mainly with respect to their electrical properties, since their optical role in the cell is of less importance (Bivour et al., 2014b). Doing this can reduce FF losses (Kobayashi et al., 2013). The n-type a-Si:H/TCO interface is less critical than the p-type a-Si:H/TCO interface. An Ohmic contact for electrons is needed, and this is relatively simple to realize in practice with sufficient doping of the n-type a-Si:H and TCO layers. Band offsets at the a-Si:H/c-Si interfaces also play an important role in carrier transport. For typical a-Si:H films, the conduction and valence band offsets are around 0.15 and 0.45 eV, respectively, roughly independent of substrate and film doping type (Fuhs et al., 2006; Korte and Schmidt, 2011; Schulze et al., 2011). Nevertheless, the valence band offset increases with the a-Si:H hydrogen content, for example, and can reach 0.6 eV (Schulze et al., 2011). A valence band offset that is too large can have a dramatic effect on hole transport, blocking carriers and reducing FF (Seif et al., 2014). Based on the lifetime measurements shown in Fig. 2.3, FF losses from recombinative processes can also be anticipated for heterojunction solar cells, and are expected to be higher for p-type cells than for n-type cells. Indeed, the minority-carrier density decreases from high to low values during an illuminated current–voltage measurement when moving from open- circuit to short-circuit conditions. The reduced lifetime at low injection in the p-type case reduces performance at maximum power point (MPP) compared to the n-type case, where the lifetime stays constant for decreasing injection (Fig. 2.3). High Voc values are not sufficient to guarantee high FF values: Even though n- and p-type heterojunction solar cells have similar Voc values, cells on p-type wafers are less efficient (Descoeudres et al., 2013). Since the minority-carrier lifetime at MPP is determined only by surface recombination (provided that high-quality c-Si wafers are used), both Voc and FF depend fundamentally on the a-Si:H/c-Si interface properties (Descoeudres et al., 2013; Reusch et al., 2013).

4.3. Jsc losses

Losses in Jsc are caused by reflection and recombination, as well as transmission if the cell is bifacial or has interdigitated back contacts. Reflection losses include front-surface reflection at the anti-reflection coating and metal Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 97 fingers, and escape reflection that results when near-bandgap light that enters the cell is imperfectly trapped in the cell and exits the front. Jsc loss due to recombination is called parasitic absorption and refers to light that is absorbed but does not result in a collected electron–hole pair because the carriers recombine (or thermalize, in the case of free-carrier absorption) during transport. Parasitic absorption can occur in the absorber itself if the diffusion length is short, in highly doped supporting layers that exhibit free-carrier absorption (e.g., TCO layers), or in supporting layers that are very defective (e.g., a-Si:H layers). Current loss analysis is conveniently simplified in silicon heterojunction solar cells because their high lifetimes and consequently long diffusion lengths mean that parasitic absorption is strictly associated with light not absorbed in the wafer. Figure 2.7 shows the external quantum efficiency (EQE) and total absorbance (1-reflectance) of a 96-μm-thick silicon heterojunction solar cell with a full silver reflector (Holman et al., 2013c). The blue shaded area indicates front-surface reflection, the purple area indicates escape reflection, the green areas indicate parasitic absorption, and the red area indicates successful charge collection. Though the current loss associated with each area is given, one must be careful: current lost is not the same as current gained if the loss mechanism is removed. For example, if infrared parasitic absorption were 2 2 removed, Jsc would increase by 1.2 mA/cm , not 2.4 mA/cm , because some of the spared photons would contribute to the escape reflection instead of to EQE.

Grid shadowing (2.8 mA/cm2 = 6.1%) 100 Front-surface reflection 90 (1.4 mA/cm2 = 3.0%) 80 Escape reflection 2 70 (1.3 mA/cm = 2.8%) 60 Blue parasitic absorption (1.5 mA/cm2 = 3.2%) 50 40 IR parasitic absorption 2 30 (2.4 mA/cm = 5.3%) 20 EQE and 1-reflectance (%) J 10 Aperture-area sc (36.7 mA/cm2 = 79.8%) 0 400 600 800 1000 1200 Wavelength (nm) Figure 2.7 EQE and total absorbance of a 96-μm-thick silicon heterojunction solar cell with a full silver reflector, showing the photons lost to each mechanism and the corresponding Jsc loss. Reproduced from Holman et al. (2013c). 98 Christophe Ballif et al.

Shadowing from the front metal grid creates the largest Jsc loss. Reducing it by narrowing the finger and busbars improves cell performance only if they are simultaneously made taller or more conductive so that their contribution to the lumped series resistance does not increase. Stencil printing is one approach to achieve higher-aspect-ratio metal lines (Zicarelli et al., 2010), and copper plating is one approach to achieve more conductive lines (Hernandez et al., 2013); both are described in detail in other sections. Parasitic absorption of blue light is responsible for the next largest loss, and is due to absorption in the front a-Si:H layers (Fujiwara and Kondo, 2007; Holman et al., 2012; Jensen et al., 2002; Page et al., 2011; Psych et al., 2010; Tanaka et al., 1992). All light absorbed in the defective p-type layer is lost via fast recombination of carriers, as is 70% of light absorbed in the intrinsic layer; blue absorption in the front TCO layer is small because of its wide bandgap (Holman et al., 2012). The simplest fix is to make these layers thinner, but this comes at a price. If the intrinsic layer thickness drops below approximately 4 nm, it is no longer able to effectively passivate the surface because the electron wavefunction in the wafer pene- trates through it and Voc drops (Fujiwara and Kondo, 2007; Holman et al., 2012; Tanaka et al., 1992). If the p-type layer is made thin, FF falls, which Bivour et al. have attributed to its becoming depleted by the heavily doped n-type front TCO, reducing band bending in the wafer (Bivour et al., 2013). An alternative approach is to make the front a-Si:H layers more transparent by either widening the bandgap by alloying with oxygen or carbon (Einsele et al., 2012; Mueller et al., 2010; Seif et al., 2014), or making the bandgap indirect by growing microcrystalline silicon (μc-Si:H) (Ding et al., 2012; Olibet et al., 2010). The former approach is challenging because FF often falls as carrier transport is inhibited by larger band offsets, and the latter approach is challenging because μc-Si:H often has an amorphous incubation layer when grown on a-Si:H. Both remain active areas of research, spurred on by recent success with wide-bandgap layers by Fraunhofer ISE and Silevo (Feldman et al., 2014; Heng et al., 2013). Yet another approach is to change the device design so that all of the a-Si:H layers are at the rear (interdigitated back contacts) or at least the p-layer is at the rear (rear emitter) so that the n-layer at the front can be very thin without being depleted. The final large Jsc loss is due to infrared parasitic absorption. This light bounces around many times in the solar cell, and is absorbed parasitically by free carriers in both the front and rear TCO layers, as well as in the rear metal reflector (Holman et al., 2013a,b, 2014). As the front TCO layer must meet a stringent sheet-resistance requirement while the rear TCO layer can Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 99 be made resistive and transparent, the front TCO often dominates the losses and the only route to improvement is higher mobilities (Holman et al., 2013a). The rear metal reflector is also lossy—even for an excellent reflector like silver—for p-polarized light arriving at the back surface (with respect to the appropriate pyramid facet) above the critical angle for internal reflection. These photons are not totally internally reflected; instead they undergo attenuated internal reflection (as in a Fourier transform infrared spectroscopy measurement performed in the attenuated total reflection mode) because the evanescent wave interacts with lossy rather than perfectly transparent media (the TCO and metal layers) (Harrick and Dupre, 1966; Holman et al., 2013a; Otto, 1968). In particular, if the evanescent wave reaches the metal reflector it will excite surface plasmon polaritons that absorb the incident photon energy. This is remedied by displacing the reflector from the back of the solar cells by at least the evanescent wave’s penetration depth with a transparent rear TCO layer or a dielectric with local openings (Holman et al., 2013c). Imperfect light trapping—taken here to mean the path length enhancement of light in the wafer in the absence of absorption—is not as large a loss as is commonly thought. The random-pyramid texture that is common in all industrial monocrystalline silicon solar cells is so effective that only approximately 0.5 mA/cm2 (depending on the wafer thickness) would be gained if the wafers had perfectly Lambertian surfaces (Yablonovitch, 1982). Never- theless, Ingenito et al. recently demonstrated a path to come even closer to the Lambertian limit using periodic arrays of pyramids of different sizes and shapes at the front and rear surfaces (Ingenito et al., 2014). Front surface reflectance is also not a particularly large Jsc loss after encapsulation because the glass and polymer of the module introduce another, intermediate refractive index. Consequently, double-layer anti-reflection coatings are seldom used.

5. INDUSTRIALIZATION AND COMMERCIALIZATION 5.1. General status The first silicon heterojunction products were sold by Sanyo in 1997. Their current annual production capacities reached approximately 900 MW at the end of 2013 (now under Panasonic). Research activities in the field started at various institutes in 2004–2005, first with the demonstration of high Voc values (i.e., over 700 mV) (Olibet et al., 2007), and later with dem- onstrations of devices exceeding 22% efficiency (Bivour et al., 2014b; 100 Christophe Ballif et al.

Descoeudres et al., 2013; Munoz et al., 2012; Schulze et al., 2011). At the same time, the expiration of key Sanyo patents triggered renewed interest from several industrial actors. Since 2013, companies such as Silevo have marketed modules with cells based on a silicon heterojunction-like structure (with the addition of SiOx tunneling layers) and with typical cell efficiencies of over 21%. Sunpreme offers bifacial modules based on silicon heterojunction solar cells on 500 wafers. Choshu Industry Co., Ltd. (CIC) has also set up a 30 MW production line and has demonstrated 22.3% efficient, 600, full-area, rear-emitter cells. CIC also demonstrated 320 W 60-cell R&D modules (Kobayashi et al., 2013). CEA-INES is piloting a line with a nominal capacity of 18 MW (Ribeyron, 2014). The equipment maker Roth and Rau (now under Meyer Burger) has developed PECVD and sputtering tools specifically for silicon heterojunction cell mass manufacturing, and demonstrated 22.3% efficient solar cells on 600 CZ wafers. They have indicated a takt time of approximately 90 s for 56 wafers per carrier in production. Several industrial and research labs have exceeded the 21% efficiency barrier with devices greater than 100 cm2 in size, and the availability of production equipment is making the technology available for mass industrialization.

5.2. Material requirements

The strength of silicon heterojunction solar cells is their high Voc values of 725–750 mV for wafer thicknesses of 100–160 μm. To obtain such high voltages, minority-carrier bulk lifetimes in the millisecond range are required. This is typically achieved using n-type CZ material. Some groups have also reported good results with properly processed n-type quasi-mono wafers ( Jay et al., 2012, 2014). As discussed previously, n-type c-Si material is chosen for two reasons: It is generally less sensitive to metal impurities than p-type c-Si, and it is not sensitive to the boron–oxygen complex (Glunz et al., 2001; Korte and Schmidt, 2011; Schmidt and Cuevas, 1999). For p-type c-Si, good devices on FZ p-type wafers with Voc over 720 mV have also been reported (Descoeudres et al., 2013) but the devices made on CZ p-type wafers suffered from low lifetimes and possible defects in the space charge region. Several wafer manufacturers have demonstrated the capabil- ity to grow full n-type CZ ingots with bulk lifetimes in excess of several milliseconds, even after multiple charge pulling. The company MEMC has shown that continuous CZ pulling (CCZ) (Li, 2013) allows for the growth of 9-m-long ingots with lifetimes decreasing from 10 to 2 ms. Figure 2.8 shows the simulated impact of bulk lifetime on the efficiency of typical Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 101

19 Wafer resistivity Cell efficiency (%) 1 Ωcm 4 Ωcm 18 10 Ωcm

17 100 1000 10,000 Bulk lifetime (μs) Figure 2.8 Efficiency of standard silicon heterojunction cells as a function of bulk lifetime. The simulations were performed with PC1D using a model presented in Ballif et al. (2012). silicon heterojunction solar cells with thicknesses of 120 μm. The model, implemented in the one-dimensional device simulator PC1D, is described in Ballif et al. (2012), and assumes low-mobility doped layers to mimic the role of a-Si:H heterointerfaces. Above a bulk lifetime of 4 ms, there is only a marginal difference in efficiency between the three doping levels. The cells still have above 21% efficiency at 800 μs lifetime (whatever the doping) and the efficiency saturates at 22.3% for a lifetime above 4 ms. If processed correctly, the full 9-m-long CCZ ingot pulled by MEMC would lead to an average cell efficiency of around 22%. All of the values reported here were obtained with a nominal Voc of 732 mV (for a 4 ms bulk lifetime), Jsc of 37.8 mA/cm2, and FF of 79.9%. Notably, as silicon heterojunction solar cells operate close to high-injection conditions at MPP and both surfaces have full-area contacts, the base doping is less critical to the series resistance than in passivated emitter and rear contact (PERC) cells. Consequently, a larger doping variation—like that obtained in typical n-type ingot growth—can be tolerated. Figure 2.9 shows simulated cell efficiency as a function of doping for an improved device with quasi-perfect passivating contacts. The results are displayed for various bulk lifetimes and wafer thickness. For a cell with a thickness of 100 μm, a bulk lifetime of 4 ms, and a wafer resistivity of 3 Ω cm, we anticipate a nominal efficiency of 25%, a 2 Voc of 755 mV, a Jsc of 39.9 mA/cm , and a FF of 83.3%, which are in the range of the best reported results from Panasonic (Taguchi et al., 2014). The plot shows that, as expected, for quasi-perfect contacts and high 102 Christophe Ballif et al.

26.0 Bulk lifetime, wafer thickness 25.5 10 ms, 150 μm 25.0 10 ms, 100 μm

μ 24.5 1 ms, 60 m 1 ms, 100 μm 1 ms, 150 μm 24.0 Cell efficiency (%)

23.5

23.0 1012 1013 1014 1015 1016 1017 Wafer doping density (cm−3) Figure 2.9 Efficiency of state-of-the-art silicon heterojunction cells as a function of background doping and for various bulk lifetime and wafer thickness combinations. The simulations were also performed with PC1D. lifetime, the doping and thickness no longer play any role. For reduced bulk lifetimes, thinner wafers and higher base doping mitigate the losses, leading to efficiencies between 23.5% and 25.5%. Finally, depending on the initial quality of the silicon wafer, material improvements can be obtained by gettering or hydrogenation. CIC also reported the use of a thermal-donor annealing step.

5.3. Temperature coefficient and energy yield

The high Voc grants silicon heterojunctions an improved temperature coefficient (Green, 1998), and typical values of 0.2–0.3%/C at MPP have been reported (Batzner et al., 2011; Heng et al., 2013; Mishima et al., 2011; Taguchi et al., 2014). Indeed, coefficients as low as 0.1%/C (Seif et al., 2014) were reported in the range of 25–55 C. The discrepancies between the values for cells with equivalent Voc stem from the properties of the a-Si:H layers and from the temperature dependence of the FF: Transport through the heterocontacts can be improved by the temperature in some cases, leading to a stable or even increased FF with the temperature and, hence, to a more favorable temperature coefficient. Note that in such cases, however, the efficiency in standard test conditions might also be lower than for cells without activated transport. In general, the lower temperature coefficient of silicon heterojunction cells should ensure—depending on the climate—a typical energy gain of 3–5% relative to standard c-Si Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 103 diffused-junction cells with 0.45%/C. The high efficiency (reduced ther- malization and increased reflection of sub-bandgap light) also translates into a lower operating temperature of the modules by 2–3 C, giving an additional energy gain of 0.4–0.7% (Kurnik et al., 2011). As silicon heterojunction cells have no fired contacts and no second-diode current (so-called J02), they show near-perfect low-illumination behavior, giving a gain of around 1–2% compared to cells with fired contacts. This can increase energy yield by 4–8%, depending on the location and climate. For building-integrated photovoltaics (BIPV), where the operating temperature may be 15–20 C higher (Kurnik et al., 2011), an additional energy gain of 3–4% can be expected.

5.4. Metallization Several approaches to metalizing silicon heterojunction solar cells have been presented previously. For production, screen printing is simple and reliable and is therefore commonly used. However, the minimum finger width possible with screen printing is approximately 70 μm, resulting in shadowing losses of 5–8%, and a lot of paste must be used because of its comparatively high resistivity. Replacing busbars with multi-wires reduces the silver required for the finger, but the cost of the low-temperature soldering alloy might partly offset the reduced silver usage. In the long term, plating may be the best option for high efficiencies at low cost. However, the additional process complexity and moderate efficiency improvement, compared to screen printing, might delay its introduction until silver becomes significantly more expensive. The metallization scheme selected, together with the cells’ contacting layers (e.g., ITO, silver or nickel vanadium), impacts the choice of both the interconnection method (soldering or gluing) and the module packaging materials. In particular, soldering on low-temperature paste is a more delicate process than on high-temperature paste, and alternative approaches, e.g., with conductive adhesives, are also used by some companies. Notably, the packaging of silicon heterojunction modules can benefit from many of the approaches developed for thin-film technology, in which TCOs and low-temperature pastes are also present.

5.5. Tools and production technologies In its most simple form, silicon heterojunction technology requires a mix of traditional and new pieces of equipment. We describe here some of the key 104 Christophe Ballif et al. steps for solar cell manufacturing and, when relevant, some features of the tools. We note that several device configurations are possible. By nature, silicon heterojunction cells can be made with the emitter at the front or at the rear of the cell, and excellent results have been reported in both configurations (Descoeudres et al., 2013). This translates into an intrinsic bifaciality that can be employed, e.g., by printing silver lines directly onto the front and rear of the solar cells.

5.5.1 Wet chemistry As described in Section 3.1, most reported approaches rely on a standard saw-damage etch following by random-pyramid texturing (or these processes are performed simultaneously) followed by surface cleaning. Immedi- ately prior to a-Si:H deposition, the wafer is dipped in a diluted HF solution to strip the oxide. The texturing and cleaning can strongly impact the solar cell performance, requiring good control of the chemical quality. Though some cleaning recipes have been reported (Edwards et al., 2008; Page et al., 2006), most texturing and cleaning processes are proprietary.

5.5.2 a-Si:H layer deposition Several deposition techniques have been reported to provide high-quality a-Si:H passivation layers including microwave PECVD ( Jeon and Kamisako, 2009), inductively coupled plasmas (Psych et al., 2010; Xiao et al., 2012), and hot-wire CVD (Branz et al., 2008; Schuttauf et al., 2011; Schu¨ttauf et al., 2011). For devices with over 21% efficiency, the main reported techniques are based on parallel-plate capacitively coupled PECVD. Good results have been reported both at 13.56 and 40.68 MHz (Descoeudres et al., 2011). An advantage of parallel-plate PECVD is the direct transfer of knowledge gained from thin-film silicon films for, e.g., doped layer deposition. Achieving homogeneity (5–10% thickness variation) in a large-area plasma reactor (>1m2) is a problem that has been solved by the flat-panel display and thin-film silicon PV industries. Careful PECVD reactor design is required to ensure uniform gas distribution and a uniform plasma even close to the electrode edges (Howling et al., 2005). If frequencies of 40 MHz or higher are used, inhomogeneities should be compensated by “lenses” for large-area reactors (Sansonnens and Schmitt, 2003; Schmidt et al., 2004). The typical deposition rate of a-Si:H layers is in the range of 0.2–0.5 nm/s. It can be inferred that process times below 1 min are achievable for each of the four a-Si:H layers (i, n, i, p) deposited in a standard silicon heterojunction solar cell. As a first approximation, assuming 60 600 wafers per Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 105 batch, a well-designed production system should have four PECVD chambers. It should process such batches with a takt time of 90–120 s, including handling, giving a nominal annual capacity of 60–100 MW, depending on the uptime. If the samples are transported on carriers, there is the possibility of contamination of the chamber with a reused carrier. Therefore carrier and contamination management can play an important role in achieving good results. Figure 2.10 is an example of a PECVD tool with a central unit used to exchange the wafers from carriers.

5.5.3 Front and rear TCO, and rear metallization Plasma-assisted evaporation (ion-beam plating) (Kobayashi et al., 2012, 2013) and sputtering (Batzner et al., 2011; Descoeudres et al., 2013; Harrison et al., 2011; Hernandez et al., 2013; Lachenal et al., 2010) have been reported as effective methods for depositing the front and rear TCO layers. Sputtering is a well-established method, and the use of a rotary magnetron allows for effective target usage (>70%). With open carriers, the front and rear sides of wafers can be coated in the same run. Most reports have focused on the use of ITO as a TCO, even though ZnO has also been used

Figure 2.10 Example of a PECVD system with a lateral processing chamber. Each chamber has a dedicated carrier. The system has a nominal throughput of 2400 600 wafers per hour. Image courtesy of Meyer Burger/Roth and Rau. 106 Christophe Ballif et al. in silicon heterojunctions (Maydell et al., 2006; Munoz et al., 2011) as both the front and rear TCO. ZnO deposited by low-pressure CVD was also reported (Choong et al., 2010). High-efficiency devices require careful optimization of the transparency of the TCO (e.g., with oxygen control) and proper control of the surface dopant density to ensure good contact with subsequent metallization (Barraud et al., 2013). The TCO at the rear also acts as a natural rear reflector. The optical properties of silicon heterojunctions have been discussed in detail in previous sections and by Holman et al. (2012, 2013a,c). In particular, if the rear TCO is strongly doped, its thickness should be minimized to avoid free-carrier absorption in it; if it is lightly doped, it should be >100 nm thick to suppress plasmon excitation in the rear metal reflector. However, in this case, the gain in Jsc and hence efficiency may be offset by the additional material and equipment costs, and implementation of ideal reflectors in industrial cells will depend on the achievable deposition rates and material costs. In bifacial and IBC cells with screen-printed silver fingers on the rear TCO, a white or transparent backsheet can be used. A white rear reflector exhibits no plasmonic loss (unlike silver) and also promotes high Jsc. However, if the silver fingers are spaced too far apart, some FF losses will occur.

5.5.4 Final processing After printing, the cells are typically cured at a temperature below 250 C for 10–20 min. This step reduces the finger and busbar line resistance, can promote a mobility (and thus transparency) increase in the TCO (Figure 2.4), and anneals the defects induced in the a-Si:H layers by the sputtering process (Demaurex et al., 2012). After curing, measurement of cell current–voltage characteristics requires a different setup than that used for diffused-junction cells. Because of the high Voc and associated diffusion capacitance (Beljean et al., 2011; Feretti et al., 2013), an illumination length in excess of several hundred milliseconds is preferred, which is too long for the flashers usually used in production lines. One approach to properly measure a cell is to use a flasher to measure Jsc only, because transient effects are limited in short-circuit conditions. To subsequently sweep the current–voltage curve, continuous (and often monochromatic) LED lighting is used with an intensity set to provide the same Jsc. A sweep time of 150–700 ms can be required, depending on the device and on the specific measurement procedures. Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 107

5.6. Production costs The production costs for silicon heterojunction solar cells can be assessed relatively simply. High-quality n-type CZ material can require some adap- tation of the pulling process (slower speed) and good-quality feedstock. The combination of thin wafers (sawn to 140 μm in the foreseeable future with the potential for 120 and 100 μm farther out) and diamond sawing can lead to competitive n-type wafer pricing comparable to standard p-type CZ material. Compared to a diffused-junction production line, the PECVD and sputtering tools are different, replacing the diffusion furnace, dielectric (e.g., silicon nitride) coating tool, rear-side etching and edge isolation sta- tions, and rear-side printer. In a reasonable time frame, the cost for 1 m2 PECVD reactors and TCO or metal coating tools should come down to €1 to 2 million. Such tools have been heavily developed for the glass-coating industry, e.g., for flat-panel displays. Thus, in terms of process equipment capital expenditures, silicon heterojunction production lines should be able to approach diffused-junction lines that include selective emitters or rear- side passivation with local openings. The major production costs for silicon heterojunction technology include the depreciation of the equipment, the targets (especially ITO and silver), and the material costs of the modules. Table 2.6 gives an estimate of the cost of the transformation of wafers to modules. This model assumes a monocrystalline 600 wafer (240 cm2), a multi-wire interconnection scheme, and cell production costs that include 40 mg of silver paste and 80 nm of ITO on both sides.

Table 2.6 Estimated direct manufacturing costs for a typical 500–1000 MW cell and module manufacturing line

Type of cost Cost in €cts/Wp Wafer 14–18 Depreciation of equipment and buildings 4–7 Direct personnel costs 1–2 Materials and operation, cells 4–6 Materials and operation, modules 12–16 Electricity 0.2–0.4 Total manufacturing costs 35–49

A high yield of 95% is assumed for the cell and module lines and the average cell efficiency is assumed to be 22%. The exact numbers depend on location, the specific choice of tools, and the metallization approach. The depreciation period is assumed to be six years for the production lines and 20 years for the buildings and facilities. 108 Christophe Ballif et al.

Considering the target module efficiency and reduced balance-of-system costs, together with the expected high energy yield, an ultra-low levelized cost of electricity (LCOE) is anticipated, typically below 4€cts/kWh in a sunny country. With potential efficiency improvements, replacement of ITO with indium-free materials, and a further decrease in wafer and module material costs, a further decrease of 10% can be expected in the cost of electricity production.

6. FUTURE DIRECTIONS AND OUTLOOK

Table 2.7 summarizes some of the strengths and weaknesses of silicon heterojunction technology. In principle, the limited number of cell processing steps and the near-infinite possibilities offered by the plasma processes used to make the passivating layers and electron and hole collectors offer the prospect of a simple and low-cost production technology that yields high-efficiency cells. 320 W 60-cell modules have already been demonstrated (Kobayashi et al., 2013) and commercial modules with 19.4% efficiency are commercially available (Panasonic, 2014b). Several companies are entering production with tens-of-MW lines and many more are assessing the technology using pilot or advanced R&D tools. Silicon heterojunction technology has the possibility to rival the more standard PERC (Blakers et al., 1989) or PERL (Green, 1991) approach in terms of manufacturing costs, but with an upside potential for high efficiency. This, coupled with their high energy yield, should allow silicon heterojunction modules to achieve reduced LCOE in solar farms. For space- constrained installations, their higher efficiency translates into higher selling prices, which should enable companies to generate a higher return on investment. Silicon heterojunction modules, with their low-temperature coefficient, could also earn a leading role in the built environment if the BIPV market realizes its long-awaited growth potential. Most of the perceived challenges, such as unstable processes for contact formation or the need for high-quality surface cleaning, can be turned into a competitive advantage in the form of proprietary processes that are not easy to copy. However, a real challenge for silicon heterojunction cells is to enter into the market with high enough volume to surpass the existing players. Most companies prefer to continue with well-established technologies and favor upgrades of existing standard, diffused-junction lines, which can be depre- ciated faster and do not have the risk of a technology step. With the clear potential for 60-cell modules with a power of 320 W (or more), silicon heterojunction cells, once established, may force companies to adopt Amorphous Silicon/Crystalline Silicon Heterojunction Solar Cells 109

Table 2.7 Some of the key features of silicon heterojunction technologies Strengths Challenges Opportunities High efficiency in Requires high-quality Less sensitive to base production, potentially up ingots (with doping level, ideal for thin to >24% correspondingly higher wafers cost/kg) Intrinsic high energy yield, Low levelized cost of thanks to temperature energy, well suited to BIPV coefficient and low- sector illumination behavior Few process steps Requires good process Proprietary production control processes/differentiation Low-temperature Requires high-quality Proprietary production processes, no stress surface preparation processes Use of established low cost/ Requires quality Potential for low capital m2 coating technology production tools expenditures for (homogeneity, production equipment repeatability, control) Compatible with thin Maintaining high wafers and diamond sawing production yield with thin wafers Cell design is intrinsically Basic versions require Use of ITO as a barrier for bifacial, different products ITO, which may plating, substitute ITO with minor changes increase costs with other TCOs By definition, cells are not Requires dedicated Optimized module lifetime subject to potential- encapsulation process and reliability, benefitting induced degradation from knowledge of thin- film cells Potentially low production Mass manufacturing Higher market price of costs experience still limited high-efficiency modules to a few manufacturers similar or other advanced technologies. At the cost of some modified or added steps, silicon heterojunction cells could also evolve. An attractive candidate is the IBC cell design, which nearly eliminates parasitic losses at the front while maintaining ultra-high Voc. Indeed a 25.6% cell efficiency was reported recently using such an approach, taking full advantage of the concept of passivating contacts (Panasonic, 2014a). More generally, a-Si:H-based contacts are paving the way for low-temperature and efficient passivating contacts on all cells—including diffused-junction cells. 110 Christophe Ballif et al.

ACKNOWLEDGMENTS The authors thank B. Demaurex, J. Geissbu¨hler, J. Seif, and A. Tomasi for support in manuscript preparation. We also acknowledge support from the EU-FP7 program and the Swiss Federal Office for Energy.

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Overview of Thin-Film Solar Cell Technologies

Bernhard Dimmler1 Manz AG, Reutlingen, Germany 1Corresponding author: e-mail address: [email protected]

Contents

1. Introduction 121 2. Market Shares of TF in PV 123 3. TF Device Efficiencies in Laboratory and Industry 125 4. Future Developments of TF Technologies in PV 128 References 136

1. INTRODUCTION

Cu(In,Ga)(Se,S)2 (CIGS), CdTe, and amorphous silicon (a-Si) in various structures are the most interesting thin-film (TF) materials to directly convert light into electricity. In general, solar cells based on III–V compounds (GaAs and related) with multi-junction devices show much higher efficiencies. They are commercially applied in space solar cells and in con- centrated, terrestrial devices due to their inherently large production costs. This material class is not covered in this review. CIGS and CdTe are polycrystalline, p-type compound semiconductors and are fabricated into solar cells via combination with an n-type layer, a typically nontransparent back contact and a transparent conductive oxide (TCO) as a front contact. Both materials have been explored since the late 1970s. In the meantime, they have gained maturity and since a few years are being manufactured in increasingly high volume. With the dynamic and rapid evolution of the PV market, crystalline silicon (c-Si) continues to play the dominating role as the most important PV material with an actual market share around 90%. TF CdTe and CIGS together with amorphous/microcrystalline silicon (a-Si/μc-Si) and other structures of a-Si have the inherent advantage when it comes to fabrication over large area in fully automated

Semiconductors and Semimetals, Volume 90 # 2014 Elsevier Inc. 121 ISSN 0080-8784 All rights reserved. http://dx.doi.org/10.1016/B978-0-12-388417-6.00004-0 122 Bernhard Dimmler production lines. The advantages span the full range from feedstock to the qualified module in one factory. From a technical point of view, solar cells are fabricated by deposition of the active films on square meter product area continuously in an inline mode; additionally, micropatterning mainly by laser scribing gives high flexibility to design optimal electrical behavior of the modules in application. Very small amounts of material and much less energy are required to manufacture a TF module when compared to c-Si. With increasing maturity of technology, all three materials started to enter the PV market with increasing quantities since about 2005. About 3–5 years ago, the market share of TF PV modules has been estimated to reach 25–30% by 2012 supported by a fast-growing market and shortage of adequate quantities of quality solar-grade silicon. However, due to recovery of availability of silicon, construction of huge production volumes and professional manufacturing, c-Si has gained even more market share in recent years. Nonetheless, the advantages for TFs remain and significant cost reductions have been realized. The major disadvantage for TFs are (i) the still lower efficiencies in comparison to c-Si in the large area module, (ii) low manufacturing scale, and (iii) innovations to be transferred from laboratory are lagging behind schedule. CIGS and recently CdTe too have demonstrated cell efficiencies in the laboratory well beyond 20%, but transferring results from lab to fab has been slow and ineffective in order to catch up with c-Si until now. The main lab to fab transfer challenges are relatively complex processing as well as professional equipment design and factory operation- related issues. In addition, more investments into CIGS are required to enable manufacturing maturity to be gained and to realize the superior cost structure. To date, there are few TF examples where roadmaps have been realized, i.e., First Solar for CdTe and Solar Frontier for CIGS with manufacturing capacities at the gigawatt per year scale. Since c-Si continuous to improve, efforts have to be strengthened and focused to bring CIGS and CdTe into the market place in high volumes. With their demonstrated potentials and the transfer to high-volume manufacturing, both TF materials can compete with c-Si already and surpass if the technical and cost targets are met in due time. In the long term, TFs can gain PV leadership if intensive R&D is kept ongoing in the short to medium term with effective technological transfer to industry. The module cost target to produce well below 0.40 US$/W is achievable with existing technologies. Nevertheless, disruptive innovations mainly on the material and packaging side are still not applied with potential to further reduce costs to below 0.3 US$/W. If this is executed well, TF can truly start a revolution in the PV market to Overview of Thin-Film Solar Cell Technologies 123 supersede c-Si in the long term. High-volume investments are necessary and the right technology choice has to be made as well in order to reach this goal.

2. MARKET SHARES OF TF IN PV

From 2004/2005 onward, TF materials have gained a first level of maturity at high-volume manufacturing in a cost-competitive way. Within the past 10 years, the market share of TF PV steadily increased by about 50% per year. Coupled with the rapid growth of c-Si, PV has emerged as a mature industry and became a real competitor to conventional energy sources. Up to a couple of years ago, grid parity was thought to be magic rather than realistic. Forecasts predicted this not to happen before 2015. The evolution to reach grid parity was mainly driven by regulations set at the political level to overcome the economic obstacles and pave the way for high-volume mass production; the latter being the prerequisite to enable cost reductions. How- ever, nobody foresaw this evolution toward grid parity to happen so fast. Today, PV has reached grid parity in high- (sun belt) and even medium- radiation (e.g., middle Europe, Japan, China, or USA) regions and is already cheaper than energy produced via conventional energy sources. The evolution of PV in the last 10 years shows mean growth rates around 50% per year; the increasing market is shown in Fig. 3.1 on the basis of shipments in the market, including estimations for the next years.

Figure 3.1 Worldwide shipments of PV modules over the last decade with forecasts until 2016 (2014–2016: blue (dark gray in the print version), conservative estimate; green (gray in the print version), accelerated estimate). Data derived from EPIA (2014); SPV Market Research (2014); various published market surveys, industry announcement, press releases, and other publicly available information; data from personal experience and private communications of the author in contact with experts in thin films worldwide. 124 Bernhard Dimmler

While over the past decades mono- and multi-crystalline silicon were the first and most advanced materials, TFs gained more and more maturity and entered the market in higher volumes in the beginning of the last decade. At that time, it was predicted that TF PV was ready for massive upscaling to high-volume manufacturing and to rapidly gain market share. The window of opportunity was aided by a fast-growing market and the shortage of solar- grade silicon which hindered c-Si to follow the learning curve. Several companies developed and offered turnkey solutions for high-volume production of a-Si/μc-Si. Within about 2 years’ time, approximately 30–40 factories were constructed, each with a starting capacity between 30 and 60 MWp/a. The scale-up was based on existing TF equipment for the flat panel display (FPD) industry. Based on perceived synergies with the FPD business as well as the promised quality and cost targets for these factories, the estimated installed capacity for TF Si was expected to lead the TF market. Unfortunately, none of the turnkey suppliers could fulfill their promised efficiency, productivity and cost targets, and most investments got lost. At the same time, a rapid increase in solar-grade silicon supply and big installations of new production facilities mainly in Asia for c-Si induced a dramatic fall of prices for c-Si modules which in consequence made production of a-Si modules in these turnkey factories noncompetitive. However, CdTe could meet and even surpass expectations. That was realized by one company, First Solar, which had a solid technical and economic foundation coupled with the right strategic and economic decisions at that time. In 2006, First Solar already decided to rapidly scale to a capacity greater than 2 GWp/a until 2011. At that time, First Solar also had cost leadership over all competing PV technologies and continues to do so to date. However, the margin has been eroded due to disadvantageous product attributes—relatively low efficiency and small substrate size. CIGS is the most complex material to produce also gained maturity in the past decade and is also there to meet the competitive cost expectations in large volumes. The Japanese company Solar Frontier is running production in the gigawatt range adding actually more capacity. A few other CIGS companies are running production lines in the 100 MWp/a range. Shipment data for TF modules are shown in Fig. 3.2. Comparing the data in Figs. 3.1 and 3.2, the share of TF materials in the worldwide PV market was peaking at about 17% in 2009 but—even with further shipment growth until 2011—lost pace with c-Si in the following years again. All together, the amount of TF modules installed was 1.3 GW in 2009, 2.3 GW in 2010, and 3.2 GW in 2011. That represents an increase from year Overview of Thin-Film Solar Cell Technologies 125

Figure 3.2 PV module shipments for the main TF materials from 2004 until 2012 (EPIA, 2014; SPV Market Research, 2014; various published market surveys, industry announcement, press releases, and other publicly available information; data from personal experience and private communications of the author in contact with experts in thin films worldwide). to year by +78% in 2009, +72% in 2010, and +33% in 2011. From 2012/2013 onward, Solar Frontier and others brought in a total of about 1.5 GWp/a of new CIGS capacity which is not yet included in Fig. 3.2. Due to facts of shut down most of the turn key lines and a few remaining still operated capacities the a-Si/μc-Si share has drastically decreased and is estimated to be in the range of well below 1 GWp/a of residual and operable capacity. A further increase of a-Si market share is not expected as efficiency and cost targets could not been met in a competitive way. Actually, it even seems to happen that a-Si is leaving the market. In the years of c-Si overcapacity (2011–2013), due to a lagging market demand and insecurities of market development, rather no investments for new capacities for CdTe or CIGS were made. Therefore, the share of TF even decreased to slightly less than 10% in 2013. A recovery is expected as all TF companies are improving their productivity and module quality on the one side and due to technical improvements in the factories on the other side. At present, the technical and cost roadmaps of TF companies are modified and shifted to better values with respect to intensified R&D efforts and recent improvements in quality with laboratory progress.

3. TF DEVICE EFFICIENCIES IN LABORATORY AND INDUSTRY

Intensive R&D work has been done on TFs within the last decades. Due to the fact that the number of experts working in the field has been 126 Bernhard Dimmler much less than for c-Si which additionally always profited from huge improvements of knowledge of c-Si technology in the semiconductor area. More initiatives in TF have been started worldwide as material qualities were showing competitiveness with c-Si in recent years. The evolution of champion efficiencies for TFs small cells is depicted in Fig. 3.3. In the last 2 years, several groups succeeded to surpass the long-lasting record of 20.4% efficiency for a small multi-crystalline silicon solar cell (Green et al., 2014). This is also shown for comparison in Table 3.1. Addi- tionally, best small-sized modules (monolithically integrated minimodules which size from a few square centimeters up to 3030 cm2) and full-size product area modules in the range of 0.7 to bigger than 1 m2 are included in Table 3.1. Several research institutes and companies are approaching 21% cell efficiency with CIGS and CdTe. What is remarkable is that more and more company research led to highest cell and module quality as these companies are investing in internal R&D to accelerate innovations in order to stay competitive in the future. Additionally, these companies are highly interested and effective to transfer these champion lab results to module size and production site. What is also remarkable that in CIGS there are still different deposition techniques for the light-absorbing layer applied with similar device qualities. A one-step inline approach with simultaneous thermal evaporation of the single elements is still competing with the so-called

Figure 3.3 Evolution of record efficiencies for TF small cells in laboratory (various published market surveys, industry announcement, press releases, and other publicly available information; data from personal experience and private communications of the author in contact with experts in thin films worldwide). Table 3.1 Summary of champion thin-film small-cell, minimodule, and full-size modules ranging from 0.7 to 1.2 m2 Samsung ZSW Solibro LG EI stab. Solar Frontier SDI Stuttgart NREL Hanergy EMPA HZB First Solar eff. CIGSS CIGSS CIGS CIGS CIGS CIGS CIGS CdTe Thin Si Efficiency 20.9 20.8 20.6 20.5 20.4 19.4 20.4 14.4 Lab cell 0.5 cm2 0.5 cm2 1cm2 0.5 cm2 Efficiency 17.8 18.7 Minimodule 3030 44 Efficiency 14.6 15.7 Manz 14.2 17.0 total area 14.6 module Remark Sequential Sequential Coevap. Coevap. Coevap. Coevap. Coevap. Sublimation PECVD Cd-free stab. eff, stabilized efficiency. Various published market surveys, industry announcement, press releases, and other publicly available information; data from personal experience and private communications of the author in contact with experts in thin films worldwide. Data from priv. comm. and press releases. 128 Bernhard Dimmler sequential processing separating material deposition and crystallization in Se/S atmospheres. This fact shows that it is possible to fabricate CIGS in different ways with similar cell qualities; therefore, there seems to be still a lot of room to improve in the direction of 25% cell efficiency which is expected to be reached with CIGS-based devices in the next 5 years. In Tables 3.2 and 3.3, the main companies operating production volumes larger than 30 MWp/a with CIGS and CdTe are listed. Table 3.2 shows the list of European companies, whereas Table 3.3 adds CIGS companies from the rest of the world. Solar Frontier is the biggest producer in the CIGS field. It has reached about 1 GWp/a capacity, a level which is necessary to reduce cost via scaling benefits. Other companies are already in the 10–100 MW range, partially with new and innovative deposition techniques and materials. After a few start-up companies in CdTe in the USA (e.g., PrimeStar, Abound Solar/GE) did not successfully reach a larger scale of operation in recent years, there remains only First Solar as the biggest TF module producing company worldwide and—at much lower level of capacity— Calyxo, a former Q-Cells company, which is located in Germany. The main data are collected in Table 3.4. First Solar is by far the most advanced TF company and had in 2010/2011 a maximum capacity of 2.7 GWp/a spread over several factories worldwide. Due in part to the large-scale volume, First Solar had the lowest manufacturing costs, but this advantage has been eroded due to increasing pressure from Asian c-Si manufacturers and upcoming CIGS players. Actu- ally, they are running about 2.2 MWp/a after closing their European factory located in Frankfurt Oder/Germany. Recently, they succeeded to show—as several CIGS companies—that cell efficiencies well beyond 20% is also possible with CdTe as absorber material. Based on intensive internal R&D, they have impressive efficiency and cost roadmaps for the coming years; if realized, they will stay well competitive in the market. First Solar today is producing CdTe modules in factories in Perrysburg/ Ohio/USA and Malaysia.

4. FUTURE DEVELOPMENTS OF TF TECHNOLOGIES IN PV

As stated above, most of the TF companies are very active doing internal R&D to stay competitive also in the next 10–20 years’ time frame. Addi- tionally, several public institutes are making research on a high know-how and quality level. As competition in the PV market is rather dynamic, Table 3.2 List of European-based companies working in pilot operation or already manufacturing in larger volumes serving the market with high-quality CIGS modules Current Champion nameplate Champion Deposition product Capacity Product minimodule Manufacturer technology ta (%) (MW/a) Product size apa (%) Remarks Avancis (G) 2-Step: Sputter +Se- 14.9 apa 30+100 Glass–glass CdS (Cd-free) Lab: Temporarily ceased JV(SK) evap.+RTP-cryst./ approx. 100 6641587 mm2 ¼1.05 m2 3030 cm2 operation since Hyundai H2S 13.6 taa 16.6 Cd-free mid-2013 In takeover by CNBM (CN) Never in full production Solibro (G) 1-Stage 15.6 apa 120 Glass–glass CdS buffer 44cm2 Overtaken by coevaporation inline approx. 7901190 cm2 ¼0.88 m2 18.7 Hanergy/China 14.3 taa Not in full operationa Announced 300 MW/a Bosch Solar 2-Step: Sputter 15.1 apa Approx. Glass–glass CdS buffer Pilot operation CISTech (G) +Se/S+H2S/H2Se approx. 5–10a 5001200 cm2 ¼0.6 m2 13.4 taa Solarion (G) 2-Step: Seion beam Approx. PI web, glass–glass flexible Pilot operation small supported 1–3a scale Product development Continued Table 3.2 List of European-based companies working in pilot operation or already manufacturing in larger volumes serving the market with high-quality CIGS modules—cont'd Current Champion nameplate Champion Deposition product Capacity Product minimodule Manufacturer technology ta (%) (MW/a) Product size apa (%) Remarks Manz 1-Stage 15.9 apa 30!8 Glass–glass CdS buffer Operation in (Wu¨rth) (G) coevaporation inline 14.6 ta (innoline) 6001200 cm2 ¼0.72 m2 innovation mode Only turn key supplier worldwide Flisom (CH) 3-Stage Approx. PI web Pilot operation small coevaporation 0.5a scale Nexcis (F) 2-Step electroplate– 13.2 apa Approx. Glass–glass 14.2 on Pilot operation small Selenization 0.5a submodule scale aEstimation of the author. apa, aperture area; ta, total area; PI, polyimide as flexible substrate. Various published market surveys, industry announcement, press releases and other publicly available information; data from personal experience and private communications of the author in contact with experts in thin films worldwide. Table 3.3 List of US- and Asian-based companies working in pilot operation or already manufacturing in larger volumes serving the market with high-quality CIGS modules Current Champion nameplate Champion Deposition product Capacity Product minimodule Manufacturer technology ta (%) (MW/a) Product size apa (%) Remarks Solar 2-Step: Sputter 14.6 1000 Glass–glass Cd-free buffer 17.8 on Announced for 2015: Frontier ( J) +H2Se/H2S- 9771257 mm2 ¼1.22 m2 3030 cm2 new factory with crystallization 12.6 with +150 MW/a CZTS device STION 2-Step: Sputter 14.5 75+5 Glass–glass CdS buffer 23.2 Tandem (USA) +H2Se/H2S- 15.7 pilota 6561656 mm2 ¼1.08 m2 device TSMC (TA) crystallization 100 MiaSole´ Reactive sputter 15.7 apa 75a ss web, glass–glass, buffer Overtaken by Hanergy/ (USA) CdS sputter China 6651611 mm2 ¼1.07 m2 Glass–glass module, prototype flexible Samsung 2-Step: Sputter 15.7 1a Glass–glass CIGS process not clear SDI (SK) +H2Se/H2S- 9001600 mm2 ¼1.44 m2 small pilot operation crystallization Global Solar 3-Stage 13 40a ss web, glass–glass Overtaken by Hanergy/ Energy coevaporation China (USA) SIVA Power 3-Stage Glass Lab cell 18.8% Start up in CIGS, small, (USA) coevaporation rather lab aEstimation of the author. apa, aperture area; ss, stainless steel flexible substrate. Various published market surveys, industry announcement, press releases, and other publicly available information; data from personal experience and private communications of the author in contact with experts in thin films worldwide. Data from press releases of the companies, status May 2014. Table 3.4 Companies producing CdTe in larger amounts Current Champion nameplate Deposition product ta Capacity Product Champion lab cell Manufacturer technology (%) (MW/a) Product size active area (%) Remarks First Solar Medium vacuum 17.0 Approx. Glass–glass 20.4 Global player (USA) thermal evaporation 2200 6001.200 mm2 ¼0.72 m2 Production lines in USA and Malaysia Calyxo (G) Atmospheric 13.4 Approx. Glass–glass thermal evaporation 80 6001.200 mm2 ¼0.72 m2 apa, aperture area. Various published market surveys, industry announcement, press releases and other publicly available information; data from personal experience and private communications of the author in contact with experts in thin films worldwide. Data from press releases of the companies, status May 2014. Overview of Thin-Film Solar Cell Technologies 133 technical, and economic progress that has to be realized mainly relies on quality and productivity features. Overall, fabrication cost of the module is the major prerequisite for being competitive. Finally, the cost of electricity produced (LCoE, levelized cost of electricity) is the main driver of competitiveness in the ambient of electricity production also with conventional sources and other renewable energies. In Table 3.5, the necessary bench- marks for the numbers to reach are collected. The depicted data are estimations of the author which are realistic and necessary to reach in order to stay competitive in the electricity market. In the long-term champion cell efficiencies will go to 25% and above whereas mean module efficiency can approach 20% in the long term. Min- imum total manufacturing cost (CoO) is actually about 0.4€/Wp for medium-sized production lines of about 150 MWp/a capacity and can be reduced by another 30–40% with actually known techniques and materials; finally, the costs are limited by today’s module design which is based on glass laminates. Together with a cost decrease to be expected also on the BOS side system cost can go down to a minimum of about 0.5€/Wp with TF modules. PV systems based on these estimations can come down to an LCoE of about 5–6€/kWh in middle European climate conditions. These numbers are optimistic and based on keeping and even enhancing high-level R&D around the main companies. Collaborative R&D actions should be decided even together with competitors joining forces to overcome low market presence of today in competition with c-Si technology. As bill of materials (material cost) of TF modules is about 50% of total cost whereas capital expenditure (capex) is in the range of about 20–30%,

Table 3.5 Numbers to be reached in the short, medium, and long term for thin-film PV champion cells and modules as well as cost of ownership (CoO), total system cost, and levelized cost of electricity (LCoE) Lab cell eff. act. Av. prod. module CoO System LCoE in middle area (%) eff. ta (%) (€/W) cost (€/W) Europe (€/kWh) Short term 23 16 0.35 <18 5a Medium 25 18 0.3 0.8 6 term 10 a Long term >25 20 0.25 0.6 5 15–20 a 134 Bernhard Dimmler these cost categories are most important to work on. Other costs of personnel are about 5% and electricity about 6–8% (database Manz AG). Electricity will become an important factor due to overall sustainability of PV in the longer term. Energy payback times for TF modules—equal for all materials—are with large volume manufacturing already in the range of about 8 months for middle European climate and with further improvements in overall qualities that will remain a main advantage in comparison to c-Si. For the energy payback time of c-Si wafer technology is it not possible to come down below the range of 18–24 months due to the high energy consumption for material cleaning and crystallization. Therefore, the main aspects in order to reduce fabrication costs and the most important topics for R&D for TF in the future are mainly by reducing • material cost by • using less pure feedstock materials • increasing material utilization during deposition • reducing active layer thicknesses • minimization of interconnection losses • on-site glass production at high line capacities 500 MWp/a • replacement of glass by cheap but highly transparent webs with high barrier characteristics • capex by improvements in • throughput/tact time per unit per equipment • process and manufacturing yield • equipment availability/uptime Module area will be increased for next-generation TF PV modules up to the order of 1.5–2 m2 by next generation and standardization of equipment; throughput will be increased by a factor of 5–10 in comparison to today. Higher maturity of equipment and more professional operation will opti- mize production yields and availability. As it is clear for TF PV technology, there are still high cost reduction and quality improvement potentials not yet applied to large volume production lines; so to say that TF technology with first large volume manufacturing of only two main players as First Solar for CdTe and Solar Frontier for CIGS is just entering the learning curve which is of the order of about 20% lower than c-Si productions. The shape of the learning curve will be steeper due to implementation of upgrades in production lines in a professional way during at least the next 10–20 years. TF PV has the chance to become the main player in PV materials. In order to reach that level, some aspects to follow are listed in the following: Overview of Thin-Film Solar Cell Technologies 135

• acceleration of transfer results and know-how from lab to fab; more focused lab R&D to manufacturing needs and conditions with a closer link of laboratories to factories and setting priorities including more collaborative work to accelerate development by joining forces • topics in fundamental R&D: • better understanding of device physics, improving device quality • long-term replacement of scarce and high cost elements like indium or tellurium • improved material quality, reduced thickness, gradient, and band gap modifications (esp. with CIGS with, e.g., Ga content up) • improvement and modifications of contact and buffer layers to increase overall absorption and to enhance the useful portion of the light spectrum • optimization of contacts to reduce carrier losses • use of novel or modified materials, multispectral devices (tandem and triple band gap devices), up- and down conversion layers, improvement of optics by reducing reflectivity and parasitic light absorption in TCO layers • use of combinations with organic layers and c-Si materials • applied R&D: • reduction of interconnect and module edge areas • modification of patterning techniques, combination of laser scribing with printing techniques • other interconnect schemes, combination of scribing, printing, and bus bars • all back contact devices • other glasses than soda lime sheet glass, speciality glasses (higher soft- ening temperature, higher strength, lower thickness, and weight) • flexible and thin substrate like stainless steel or PI webs of very thin, flexible speciality glasses • high isolation barrier coatings to avoid second glass lamination • nonvacuum, atmospheric pressure deposition techniques like printing, etc. • introduction and prequalification of new inline metrology tools • modifications of devices for further reduction of temperature coefficients • improvement of device quality, sealing components, and materials to enhance stable operation beyond 30 years • environmentally sustainable materials and recycling issues 136 Bernhard Dimmler

• industrial R&D, introduction of prequalified materials and processes: • increase of module efficiency and stability, introducing results from laboratory to volume production • increase of throughput by higher deposition rates and shorter heating and cooling cycles • next-generation equipment for bigger substrate sizes • improvement of factory layouts and logistics applying, e.g., automo- tive and semiconductor concepts to increase factory yield and availability • integration of better metrology tools for quality assurance • standardization of equipment, materials, and components • minimization of energy needed for materials and production This list of topics is for sure not complete; nevertheless, it covers the main features for future improvement for competiveness in the PV market. The key and decisive factor is scaling to high volumes in order to profit from cost reductions by reduction of capital expenditure of equipment and infrastruc- ture and by purchasing high quantities of materials at lower cost. Fulfilling these tasks and aspects for continuous improvement of module qualities and manufacturing productivity to reduce cost of modules together with a reduction in BOS, TF PV materials have a good chance getting bigger in size and market share and becoming even more competitive in the long- time frame. There are still a lot of potentials to gain, but there has to be developed the right industrial strategy supported by an adequate industry policy. REFERENCES EPIA, 2014. Global Market Outlook 2014–2018. EPIA, Intersolar, Munich. Green, M.A., Emery, K., Hishikawa, Y., Warta, W., Dunlop, E.D., 2014. Solar cell efficiency tables (version 43). Prog. Photovolt. Res. Appl. 22, 1–9. Manz AG. Data from Manz internal CIGS database. Reutlingen, Germany. SPV Market Research, 2014. Photovoltaic Manufacturer Shipments: Capacity, Price & Rev- enues 2013/2014: Report SPV-Supply2, SPV Market Research, April 2014, Paula Mints. INDEX

Note: Page numbers followed by “f ” indicate figures and “t” indicate tables.

A First Solar, 128 Amphoteric defect, 78–79 large volume manufacturing, 134 a-Si:H. See Hydrogenated amorphous Solar Frontier, 124, 128 silicon (a-Si:H) US- and Asian-based companies, 131t Atmospheric pressure chemical vapor deposition (APCVD), 27 D Auger recombination, 15–16 Deep-defect recombination, 78–79 Diffused-junction solar cells B best cells, 92, 94t Back surface field (BSF), 9 cell current–voltage characteristics, 106 Boron-doped a-Si:H film, 85, 86 diffused emitter, 86–87 Building-integrated photovoltaics (BIPV), metallization, 90–92 102–103 p-doped c-Si, 82–83 production costs, 107 C temperature coefficient, 102–103 Calyxo, 128 Doped a-Si:H films Carrier recombination, silicon carrier selective contacts, 85 defect recombination, 74, 75, 78–79, 95 cross-contamination, 86 intrinsic recombination, 74, 75–76 molybdenum oxide (MoOx), 86 CdTe, 121–123 PECVD, 86 cell efficiency research, 126–128 E First Solar, 124, 128 large volume manufacturing, 134 Edge-defined film-fed growth (EFG), 8 f f manufacturer, 132t Ethylene vinyl acetate (EVA), 49, 49 ,50 Chemical surface passivation, 75–76 F CIGS. See Cu(In,Ga)(Se,S)2 (CIGS) Continuous CZ pulling (CCZ), 100–102 Field-effect passivation, 75 Crystalline silicon (c-Si) Fill factor (FF) losses, 95–96 a-Si:H/c-Si heterojunction solar cells First Solar, 124, 128 (see Silicon heterojunction solar cells) defect recombination, 75 H industrial c-Si solar cell (see Industrial Heterojunction with intrinsic thin layer crystalline Si (c-Si) solar cell) (HIT) concept, 6–7, 39–40 passivating contacts, 75–76 Honeymoon cell, 76–77 TF PV materials, 121–123 Hydrogenated amorphous silicon (a-Si:H) energy payback time, 134 as-deposited polycrystalline layers, 80–81 large volume production lines, 134 carrier recombination, 78–79 R&D work, 125–126 doped a-Si:H films turnkey solutions, 124 carrier selective contacts, 85 Cu(In,Ga)(Se,S)2 (CIGS), 121–123 cross-contamination, 86 cell efficiency research, 126–128 molybdenum oxide (MoOx), 86 European-based companies, 129t PECVD, 86

137 138 Index

Hydrogenated amorphous silicon (a-Si:H) drawbacks and challenges, 35–36 (Continued ) dual print, 36 epitaxial growth, 81 multi-busbar approaches, 39–40 honeymoon cell, 76–77 paste development, 38 H2 plasma treatments, 81–82, 82f seed-and-plate, 38–39 isothermal annealing, 79 stencil printing, 36–37, 37f mechanism, 79 fundamental loss mechanisms, 13–14, 13f n- and p-type wafers, 82–83 heterojunction emitter technology, 61 Panasonic, 77–78 IBC-HIT approach, 6–7 PECVD, 79–80, 104–105, 105f j–V characteristics, 12–13, 12f, 33–34 planar deposition technology, 76–77 n-type Si wafers, 61 solar cell efficiency, 76–78, 102–103 optical and electrical losses, 18, 19f TCO layers (see Transparent conductive output power, 12–13, 12f oxide (TCO) layers) PERC-type solar cells (see Passivated emitter and rear cell (PERC)) I PERL, 51, 52f IBC solar cells. See Interdigitated-back- PERT, 51, 52f contacted (IBC) solar cells photovoltaic (PV) Indium tin oxide (ITO), 88–89 German PV market, 7–8 absorbance spectra, 90f market share of, 8, 8f external quantum efficiency and solar modules, yearly production/ 1-reflectance spectra, 90f shipment of, 7–8, 7f thickness, mobility and free-electron in space/terrestrial applications, 4–5 density, 89f wafer and module costs, 9, 10f Industrial crystalline Si (c-Si) solar cell radiative recombination, 14–15 alkaline wafer texturing, 35 rear contacted solar cells, 61 Auger recombination, 15–16 saturation current density, 18 band diagram, 10–12, 11f SRH recombination, 16–17 basic structure of, 9–10, 9f surface recombination, 17 complex processing scheme, 5 two-diode model, 14, 14f cost reduction, 5–7 wafer sawing, 34–35 1-diode model, 12 Interdigitated-back-contacted (IBC) solar efficiency of, 5 cells, 6–7, 92–95, 94t emitter formation Intrinsic recombination, 74, 75–76 homogeneous emitters, 41–42, 42f selective emitters (see Selective emitter L (SE)) Laser fired contacts (LFC), 54 firing through SiNx:H process edge isolation, 25 M inline/batch processing, 19 Molybdenum oxide (MoOx), 86 P diffusion, 22–25, 23f Monocrystalline Si (mono-Si), 6 screen-printing metallization Multicrystalline Si (mc-Si), 6 (see Screen printing) SiNx:H deposition, 25–27 O wafer washing, texturization, and Open-circuit voltage (Voc) losses, 95 cleaning, 20–22, 21f front contact metallization P dispensing, 37–38 Passivated emitter and rear cell (PERC) double print, 36 boron–oxygen related degradation Index 139

metastable, 57 full silver reflector, 97, 97f regenerated state, 58–59, 59f infrared parasitic absorption, 98–99 strategies, 57–58 Lambertian surfaces, 99 dielectric rear side passivation recombination losses, 96–97 Al2O3 layers, 53 reflection losses, 96–97 Al2O3/SiNx:H stacks, 53–54 Silicon SiO2/SiNx:H stacks, 54 carrier recombination efficiencies of, 59–60 defect recombination, 74, 75, 78–79, 95 local rear contact formation intrinsic recombination, 74, 75–76 contact patterns and void formation, emitter diffusion process, 74 55–56 Silicon dangling bond, 78–79 etching paste, 55 Silicon heterojunction solar cells interconnection ribbons, soldering of, a-Si:H, surface passivation 56–57 as-deposited polycrystalline layers, laser fired contacts, 54 80–81 local ablation, 55 carrier recombination, 78–79 Phosphor silicate glass (PSG), 22 epitaxial growth, 81 Phosphorus-doped a-Si:H film, 85, 86 honeymoon cell, 76–77 Plasma-enhanced chemical vapor deposition H2 plasma treatments, 81–82, 82f (PECVD), 26–27, 79–80, 86, isothermal annealing, 79 104–105, 105f mechanism, 79 n- and p-type wafers, 82–83 S Panasonic, 77–78 Screen printing, 28f, 90–92, 103 PECVD, 79–80, 104–105, 105f co-firing step, 32–33, 33f planar deposition technology, 76–77 front side metallization, 28–29, 30f solar cell efficiency, 76–78, 102–103 inline process, 27–28 band diagram, 77–78, 78f lifetime of, 27–28 CEAINES, 99–100 rear side metallization, 29–32, 31f cell current–voltage characteristics, 106 Seed-and-plate approach, 38–39 challenges, 108–109, 109t Selective emitter (SE) chemical surface passivation, 76 buried contact approach, 43 Choshu Industry, 99–100 doped Si inks, 44 curing process, 106 efficiency potential of, 50–51 current annual production capacities, etch-back process, 45–48, 45f,46f,47f 99–100 EVA and silicone, encapsulation, 49–50, deposited emitter, 86–87 49f,50f device configurations, 103–104 ion implantation process, 44–45 efficiency vs. bulk lifetime, 100–102, 101f laser doping electron and hole collectors, 85–86 laser chemical processing, 48 energy yield, 102–103 NiAg light-induced plating, 48 fill factor (FF) losses, 95–96 P-glass, 48 front TCO layers, 86–87, 105–106 and plating, 48 IBC solar cells, 92–95, 94t oxide mask process, 44, 44f low-temperature copper paste, 91 restrictions, 43 low-temperature soldering, 103 Shockley–Read–Hall recombination material requirements, 100–102 statistics, 16–17, 78–79 minority-carrier effective lifetimes, 84f Short-circuit current density ( Jsc) losses monocrystalline silicon CZ wafer, 83–85 blue absorption, 98 multi-wire contacting scheme, 91, 91f 140 Index

Silicon heterojunction solar cells (Continued ) T V open-circuit voltage ( oc) losses, 95 Thin-film photovoltaic (TF PV) modules opportunities, 109t applied R&D, 135 parallel-plate PECVD, 104 Calyxo, 128 PC1D simulations, 100–102, 101f, 102f capital expenditure, 133–134 plated contacts, 92, 103 CdTe (see CdTe) production costs, 107–108 cell efficiency, 125–128, 126f random-pyramid texturing, 83–85 CIGS (see Cu(In,Ga)(Se,S)2 (CIGS)) rear TCO layers, 87–88, 105–106 collaborative R&D actions, 133 research activities, 99–100 cost of ownership (CoO), 133, 133t Sanyo patents, 99–100 c-Si (see Crystalline silicon (c-Si)) screen printing, 90–92, 103 energy payback time, 134 J short-circuit current density ( sc) losses First Solar, 128 blue absorption, 98 full-size module champion, 127t full silver reflector, 97, 97f fundamental R&D, 135 infrared parasitic absorption, 98–99 grid parity, 123 Lambertian surfaces, 99 industrial R&D, 136 recombination losses, 96–97 large volume production lines, 134 reflection losses, 96–97 levelized cost of electricity, 128–133, Silevo, 99–100 133t strengths, 108–109, 109t long-term champion companies, cost Sunpreme, 99–100 estimations of, 128–133 temperature coefficient, 102–103 market share, 121–125 thin-film deposition technology, 76 material cost, 133–134 vs. silicon homojunction solar cell minimodule champion, 127t technology, 76–77 module area and throughput, 134 wafer cleaning and texturing, 85, 104 shipped modules vs. year, 125f wafer doping density vs. efficiency, small-cell module champion, 127t 100–102, 102f Solar Frontier, 128 wet chemistry, 104 total system cost, 133, 133t Solar Frontier, 128 turnkey solutions, 124 Stencil printing, 36–37, 37f worldwide shipments, 123f Surface passivation, silicon heterojunction Transparent conductive oxide (TCO) layers solar cells front layers, 86–87, 105–106 as-deposited polycrystalline layers, 80–81 ITO and IO:H layer, 88–89 carrier recombination, 78–79 absorbance spectra, 90f epitaxial growth, 81 external quantum efficiency and honeymoon cell, 76–77 1-reflectance spectra, 90f f H2 plasma treatments, 81–82, 82 thickness, mobility and free-electron isothermal annealing, 79 density, 89f mechanism, 79 rear layers, 87–88, 105–106 n- and p-type wafers, 82–83 Two-diode model, 14, 14f Panasonic, 77–78 PECVD, 79–80, 104–105, 105f planar deposition technology, 76–77 U solar cell efficiency, 76–78, 102–103 Upgraded metallurgical grade (UMG), 5–6 CONTENTS OF VOLUMES IN THIS SERIES

Volume 1 Physics of III–V Compounds

C. Hilsum, Some Key Features of III–V Compounds F. Bassani, Methods of Band Calculations Applicable to III–V Compounds E. O. Kane, The k-p Method V. L. Bonch–Bruevich, Effect of Heavy Doping on the Semiconductor Band Structure D. Long, Energy Band Structures of Mixed Crystals of III–V Compounds L. M. Roth and P. N. Argyres, Magnetic Quantum Effects S. M. Puri and T. H. Geballe, Thermomagnetic Effects in the Quantum Region W. M. Becker, Band Characteristics near Principal Minima from Magnetoresistance E. H. Putley, Freeze-Out Effects, Hot Electron Effects, and Submillimeter Photoconductivity in InSb H. Weiss, Magnetoresistance B. Ancker-Johnson, Plasma in Semiconductors and Semimetals

Volume 2 Physics of III–V Compounds

M. G. Holland, Thermal Conductivity S. I. Novkova, Thermal Expansion U. Piesbergen, Heat Capacity and Debye Temperatures G. Giesecke, Lattice Constants J. R. Drabble, Elastic Properties A. U. Mac Rae and G. W. Gobeli, Low Energy Electron Diffraction Studies R. Lee Mieher, Nuclear Magnetic Resonance B. Goldstein, Electron Paramagnetic Resonance T. S. Moss, Photoconduction in III–V Compounds E. Antoncik and J. Tauc, Quantum Efficiency of the Internal Photoelectric Effect in InSb G. W. Gobeli and I. G. Allen, Photoelectric Threshold and Work Function P. S. Pershan, Nonlinear Optics in III–V Compounds M. Gershenzon, Radiative Recombination in the III–V Compounds F. Stern, Stimulated Emission in Semiconductors

Volume 3 Optical Properties of III–V Compounds

M. Hass, Lattice Reflection W. G. Spitzer, Multiphonon Lattice Absorption D. L. Stierwalt and R. F. Potter, Emittance Studies H. R. Philipp and H. Ehrenveich, Ultraviolet Optical Properties M. Cardona, Optical Absorption Above the Fundamental Edge E. J. Johnson, Absorption Near the Fundamental Edge J. O. Dimmock, Introduction to the Theory of Exciton States in Semiconductors

141 142 Contents of Volumes in this Series

B. Lax and J. G. Mavroides, Interband Magnetooptical Effects H. Y. Fan, Effects of Free Carries on Optical Properties E. D. Palik and G. B. Wright, Free-Carrier Magnetooptical Effects R. H. Bube, Photoelectronic Analysis B. O. Seraphin and H. E. Benett, Optical Constants

Volume 4 Physics of III–V Compounds

N. A. Goryunova, A. S. Borchevskii and D. N. Tretiakov, Hardness N. N. Sirota, Heats of Formation and Temperatures and Heats of Fusion of Compounds of AIIIBV D. L. Kendall, Diffusion A. G. Chynoweth, Charge Multiplication Phenomena R. W. Keyes, The Effects of Hydrostatic Pressure on the Properties of III–V Semiconductors L. W. Aukerman, Radiation Effects N. A. Goryunova, F. P. Kesamanly, and D. N. Nasledov, Phenomena in Solid Solutions R. T. Bate, Electrical Properties of Nonuniform Crystals

Volume 5 Infrared Detectors

H. Levinstein, Characterization of Infrared Detectors P. W. Kruse, Indium Antimonide Photoconductive and Photoelectromagnetic Detectors M. B. Prince, Narrowband Self-Filtering Detectors I. Melngalis and T. C. Hannan, Single-Crystal Lead-Tin Chalcogenides D. Long and J. L. Schmidt, Mercury-Cadmium Telluride and Closely Related Alloys E. H. Putley, The Pyroelectric Detector N. B. Stevens, Radiation Thermopiles R. J. Keyes and T. M. Quist, Low Level Coherent and Incoherent Detection in the Infrared M. C. Teich, Coherent Detection in the Infrared F. R. Arams, E. W. Sard, B. J. Peyton and F. P. Pace, Infrared Heterodyne Detection with Gigahertz IF Response H. S. Sommers, Jr., Macrowave-Based Photoconductive Detector R. Sehr and R. Zuleeg, Imaging and Display

Volume 6 Injection Phenomena

M. A. Lampert and R. B. Schilling, Current Injection in Solids: The Regional Approximation Method R. Williams, Injection by Internal Photoemission A. M. Barnett, Current Filament Formation R. Baron and J. W. Mayer, Double Injection in Semiconductors W. Ruppel, The Photoconductor-Metal Contact Contents of Volumes in this Series 143

Volume 7 Application and Devices Part A J. A. Copeland and S. Knight, Applications Utilizing Bulk Negative Resistance F. A. Padovani, The Voltage-Current Characteristics of Metal-Semiconductor Contacts P. L. Hower, W. W. Hooper, B. R. Cairns, R. D. Fairman, and D. A. Tremere, The GaAs Field-Effect Transistor M. H. White, MOS Transistors G. R. Antell, Gallium Arsenide Transistors T. L. Tansley, Heterojunction Properties

Part B T. Misawa, IMPATT Diodes H. C. Okean, Tunnel Diodes R. B. Campbell and Hung-Chi Chang, Silicon Junction Carbide Devices

R. E. Enstrom, H. Kressel, and L. Krassner, High-Temperature Power Rectifiers of GaAs1ÀxPx

Volume 8 Transport and Optical Phenomena

R. J. Stirn, Band Structure and Galvanomagnetic Effects in III–V Compounds with Indirect Band Gaps R. W. Ure, Jr., Thermoelectric Effects in III–V Compounds H. Piller, Faraday Rotation H. Barry Bebb and E. W. Williams, Photoluminescence I: Theory E. W. Williams and H. Barry Bebb, Photoluminescence II: Gallium Arsenide

Volume 9 Modulation Techniques

B. O. Seraphin, Electroreflectance R. L. Aggarwal, Modulated Interband Magnetooptics D. F. Blossey and Paul Handler, Electroabsorption B. Batz, Thermal and Wavelength Modulation Spectroscopy I. Balslev, Piezooptical Effects D. E. Aspnes and N. Bottka, Electric-Field Effects on the Dielectric Function of Semiconductors and Insulators

Volume 10 Transport Phenomena

R. L. Rhode, Low-Field Electron Transport J. D. Wiley, Mobility of Holes in III–V Compounds C. M. Wolfe and G. E. Stillman, Apparent Mobility Enhancement in Inhomogeneous Crystals R. L. Petersen, The Magnetophonon Effect 144 Contents of Volumes in this Series

Volume 11 Solar Cells

H. J. Hovel, Introduction; Carrier Collection, Spectral Response, and Photocurrent; Solar Cell Electrical Characteristics; Efficiency; Thickness; Other Solar Cell Devices; Radiation Effects; Temperature and Intensity; Solar Cell Technology

Volume 12 Infrared Detectors (II)

W. L. Eiseman, J. D. Merriam, and R. F. Potter, Operational Characteristics of Infrared Photodetectors P. R. Bratt, Impurity Germanium and Silicon Infrared Detectors E. H. Putley, InSb Submillimeter Photoconductive Detectors G. E. Stillman, C. M. Wolfe, and J. O. Dimmock, Far-Infrared Photoconductivity in High Purity GaAs G. E. Stillman and C. M. Wolfe, Avalanche Photodiodes P. L. Richards, The Josephson Junction as a Detector of Microwave and Far-Infrared Radiation E. H. Putley, The Pyroelectric Detector – An Update

Volume 13 Cadmium Telluride

K. Zanio, Materials Preparations; Physics; Defects; Applications

Volume 14 Lasers, Junctions, Transport

N. Holonyak, Jr., and M. H. Lee, Photopumped III–V Semiconductor Lasers H. Kressel and J. K. Butler, Heterojunction Laser Diodes A. Van der Ziel, Space-Charge-Limited Solid-State Diodes P. J. Price, Monte Carlo Calculation of Electron Transport in Solids

Volume 15 Contacts, Junctions, Emitters

B. L. Sharma, Ohmic Contacts to III–V Compounds Semiconductors A. Nussbaum, The Theory of Semiconducting Junctions J. S. Escher, NEA Semiconductor Photoemitters

Volume 16 Defects, (HgCd)Se, (HgCd)Te

H. Kressel, The Effect of Crystal Defects on Optoelectronic Devices

C. R. Whitsett, J. G. Broerman, and C. J. Summers, Crystal Growth and Properties of Hg1Àx Cdx Se Alloys

M. H. Weiler, Magnetooptical Properties of Hg1Àx Cdx Te Alloys

P. W. Kruse and J. G. Ready, Nonlinear Optical Effects in Hg1Àx Cdx Te

Volume 17 CW Processing of Silicon and Other Semiconductors

J. F. Gibbons, Beam Processing of Silicon A. Lietoila, R. B. Gold, J. F. Gibbons, and L. A. Christel, Temperature Distributions and Solid Phase Reaction Rates Produced by Scanning CW Beams Contents of Volumes in this Series 145

A. Leitoila and J. F. Gibbons, Applications of CW Beam Processing to Ion Implanted Crystalline Silicon N. M. Johnson, Electronic Defects in CW Transient Thermal Processed Silicon K. F. Lee, T. J. Stultz, and J. F. Gibbons, Beam Recrystallized Polycrystalline Silicon: Properties, Applications, and Techniques T. Shibata, A. Wakita, T. W. Sigmon and J. F. Gibbons, Metal-Silicon Reactions and Silicide Y. I. Nissim and J. F. Gibbons, CW Beam Processing of Gallium Arsenide

Volume 18 Mercury Cadmium Telluride

P. W. Kruse, The Emergence of (Hg1Àx Cdx) Te as a Modern Infrared Sensitive Material H. E. Hirsch, S. C. Liang, and A. G. White, Preparation of High-Purity Cadmium, Mercury, and Tellurium W. F. H. Micklethwaite, The Crystal Growth of Cadmium Mercury Telluride P. E. Petersen, Auger Recombination in Mercury Cadmium Telluride R. M. Broudy and V. J. Mazurczyck, (HgCd) Te Photoconductive Detectors M. B. Reine, A. K. Soad, and T. J. Tredwell, Photovoltaic Infrared Detectors M. A. Kinch, Metal-Insulator-Semiconductor Infrared Detectors

Volume 19 Deep Levels, GaAs, Alloys, Photochemistry

G. F. Neumark and K. Kosai, Deep Levels in Wide Band-Gap III–V Semiconductors D. C. Look, The Electrical and Photoelectronic Properties of Semi-Insulating GaAs R. F. Brebrick, Ching-Hua Su, and Pok-Kai Liao, Associated Solution Model for Ga-In-Sb and Hg-Cd-Te Y. Ya. Gurevich and Y. V. Pleskon, Photoelectrochemistry of Semiconductors

Volume 20 Semi-Insulating GaAs

R. N. Thomas, H. M. Hobgood, G. W. Eldridge, D. L. Barrett, T. T. Braggins, L. B. Ta, and S. K. Wang, High-Purity LEC Growth and Direct Implantation of GaAs for Monolithic Microwave Circuits C. A. Stolte, Ion Implantation and Materials for GaAs Integrated Circuits C. G. Kirkpatrick, R. T. Chen, D. E. Holmes, P. M. Asbeck, K. R. Elliott, R. D. Fairman, and J. R. Oliver, LEC GaAs for Integrated Circuit Applications J. S. Blakemore and S. Rahimi, Models for Mid-Gap Centers in Gallium Arsenide

Volume 21 Hydrogenated Amorphous Silicon Part A J. I. Pankove, Introduction M. Hirose, Glow Discharge; Chemical Vapor Deposition Y. Uchida, di Glow Discharge T. D. Moustakas, Sputtering I. Yamada, Ionized-Cluster Beam Deposition B. A. Scott, Homogeneous Chemical Vapor Deposition 146 Contents of Volumes in this Series

F. J. Kampas, Chemical Reactions in Plasma Deposition P. A. Longeway, Plasma Kinetics H. A. Weakliem, Diagnostics of Silane Glow Discharges Using Probes and Mass Spectroscopy L. Gluttman, Relation between the Atomic and the Electronic Structures A. Chenevas-Paule, Experiment Determination of Structure S. Minomura, Pressure Effects on the Local Atomic Structure D. Adler, Defects and Density of Localized States

Part B J. I. Pankove, Introduction G. D. Cody, The Optical Absorption Edge of a-Si: H N. M. Amer and W. B. Jackson, Optical Properties of Defect States in a-Si: H P. J. Zanzucchi, The Vibrational Spectra of a-Si: H Y. Hamakawa, Electroreflectance and Electroabsorption J. S. Lannin, Raman Scattering of Amorphous Si, Ge, and Their Alloys R. A. Street, Luminescence in a-Si: H R. S. Crandall, Photoconductivity J. Tauc, Time-Resolved Spectroscopy of Electronic Relaxation Processes P. E. Vanier, IR-Induced Quenching and Enhancement of Photoconductivity and Photoluminescence H. Schade, Irradiation-Induced Metastable Effects L. Ley, Photoelectron Emission Studies

Part C J. I. Pankove, Introduction J. D. Cohen, Density of States from Junction Measurements in Hydrogenated Amorphous Silicon P. C. Taylor, Magnetic Resonance Measurements in a-Si: H K. Morigaki, Optically Detected Magnetic Resonance J. Dresner, Carrier Mobility in a-Si: H T. Tiedje, Information About Band-Tail States from Time-of-Flight Experiments A. R. Moore, Diffusion Length in Undoped a-S: H W. Beyer and J. Overhof, Doping Effects in a-Si: H H. Fritzche, Electronic Properties of Surfaces in a-Si: H C. R. Wronski, The Staebler-Wronski Effect R. J. Nemanich, Schottky Barriers on a-Si: H B. Abeles and T. Tiedje, Amorphous Semiconductor Superlattices

Part D J. I. Pankove, Introduction D. E. Carlson, Solar Cells G. A. Swartz, Closed-Form Solution of I–V Characteristic for a s-Si: H Solar Cells I. Shimizu, Electrophotography S. Ishioka, Image Pickup Tubes P. G. Lecomber and W. E. Spear, The Development of the a-Si: H Field-Effect Transistor and its Possible Applications Contents of Volumes in this Series 147

D. G. Ast, a-Si: H FET-Addressed LCD Panel S. Kaneko, Solid-State Image Sensor M. Matsumura, Charge-Coupled Devices M. A. Bosch, Optical Recording A. D’Amico and G. Fortunato, Ambient Sensors H. Kulkimoto, Amorphous Light-Emitting Devices R. J. Phelan, Jr., Fast Decorators and Modulators J. I. Pankove, Hybrid Structures P. G. LeComber, A. E. Owen, W. E. Spear, J. Hajto, and W. K. Choi, Electronic Switching in Amorphous Silicon Junction Devices

Volume 22 Lightwave Communications Technology Part A K. Nakajima, The Liquid-Phase Epitaxial Growth of InGaAsP W. T. Tsang, Molecular Beam Epitaxy for III–V Compound Semiconductors G. B. Stringfellow, Organometallic Vapor-Phase Epitaxial Growth of III–V Semiconductors G. Beuchet, Halide and Chloride Transport Vapor-Phase Deposition of InGaAsP and GaAs

M. Razeghi, Low-Pressure, Metallo-Organic Chemical Vapor Deposition of GaxIn1ÀxAsP1Ày Alloys P. M. Petroff, Defects in III–V Compound Semiconductors

Part B J. P. van der Ziel, Mode Locking of Semiconductor Lasers K. Y. Lau and A. Yariv, High-Frequency Current Modulation of Semiconductor Injection Lasers C. H. Henry, Special Properties of Semi Conductor Lasers Y. Suematsu, K. Kishino, S. Arai, and F. Koyama, Dynamic Single-Mode Semiconductor Lasers with a Distributed Reflector W. T. Tsang, The Cleaved-Coupled-Cavity (C3) Laser

Part C R. J. Nelson and N. K. Dutta, Review of InGaAsP InP Laser Structures and Comparison of Their Performance N. Chinone and M. Nakamura, Mode-Stabilized Semiconductor Lasers for 0.7–0.8- and 1.1–1.6-μm Regions Y. Horikoshi, Semiconductor Lasers with Wavelengths Exceeding 2 μm B. A. Dean and M. Dixon, The Functional Reliability of Semiconductor Lasers as Optical Transmitters R. H. Saul, T. P. Lee, and C. A. Burus, Light-Emitting Device Design C. L. Zipfel, Light-Emitting Diode-Reliability T. P. Lee and T. Li, LED-Based Multimode Lightwave Systems K. Ogawa, Semiconductor Noise-Mode Partition Noise

Part D F. Capasso, The Physics of Avalanche Photodiodes T. P. Pearsall and M. A. Pollack, Compound Semiconductor Photodiodes 148 Contents of Volumes in this Series

T. Kaneda, Silicon and Germanium Avalanche Photodiodes S. R. Forrest, Sensitivity of Avalanche Photodetector Receivers for High-Bit-Rate Long-Wavelength Optical Communication Systems J. C. Campbell, Phototransistors for Lightwave Communications

Part E S. Wang, Principles and Characteristics of Integrable Active and Passive Optical Devices S. Margalit and A. Yariv, Integrated Electronic and Photonic Devices T. Mukai, A. Yamamoto, and T. Kimura, Optical Amplification by Semiconductor Lasers

Volume 23 Pulsed Laser Processing of Semiconductors

R. F. Wood, C. W. White and R. T. Young, Laser Processing of Semiconductors: An Overview C. W. White, Segregation, Solute Trapping and Supersaturated Alloys G. E. Jellison, Jr., Optical and Electrical Properties of Pulsed Laser-Annealed Silicon R. F. Wood and G. E. Jellison, Jr., Melting Model of Pulsed Laser Processing R. F. Wood and F. W. Young, Jr., Nonequilibrium Solidification Following Pulsed Laser Melting D. H. Lawndes and G. E. Jellison, Jr., Time-Resolved Measurement During Pulsed Laser Irradiation of Silicon D. M. Zebner, Surface Studies of Pulsed Laser Irradiated Semiconductors D. H. Lowndes, Pulsed Beam Processing of Gallium Arsenide

R. B. James, Pulsed CO2 Laser Annealing of Semiconductors R. T. Young and R. F. Wood, Applications of Pulsed Laser Processing

Volume 24 Applications of Multiquantum Wells, Selective Doping, and Superlattices

C. Weisbuch, Fundamental Properties of III–V Semiconductor Two-Dimensional Quantized Structures: The Basis for Optical and Electronic Device Applications H. Morkoc¸and H. Unlu, Factors Affecting the Performance of (Al,Ga)As/GaAs and (Al,Ga)As/InGaAs Modulation-Doped Field-Effect Transistors: Microwave and Digital Applications N. T. Linh, Two-Dimensional Electron Gas FETs: Microwave Applications M. Abe et al., Ultra-High-Speed HEMT Integrated Circuits D. S. Chemla, D. A. B. Miller and P. W. Smith, Nonlinear Optical Properties of Multiple Quantum Well Structures for Optical Signal Processing F. Capasso, Graded-Gap and Superlattice Devices by Band-Gap Engineering W. T. Tsang, Quantum Confinement Heterostructure Semiconductor Lasers G. C. Osbourn et al., Principles and Applications of Semiconductor Strained-Layer Superlattices

Volume 25 Diluted Magnetic Semiconductors

W. Giriat and J. K. Furdyna, Crystal Structure, Composition, and Materials Preparation of Diluted Magnetic Semiconductors Contents of Volumes in this Series 149

W. M. Becker II , Band Structure and Optical Properties of Wide-Gap A1ÀxMnxBIV Alloys at Zero Magnetic Field S. Oseroff and P. H. Keesom, Magnetic Properties: Macroscopic Studies T. Giebultowicz and T. M. Holden, Neutron Scattering Studies of the Magnetic Structure and Dynamics of Diluted Magnetic Semiconductors J. Kossut, Band Structure and Quantum Transport Phenomena in Narrow-Gap Diluted Magnetic Semiconductors C. Riquaux, Magnetooptical Properties of Large-Gap Diluted Magnetic Semiconductors J. A. Gaj, Magnetooptical Properties of Large-Gap Diluted Magnetic Semiconductors J. Mycielski, Shallow Acceptors in Diluted Magnetic Semiconductors: Splitting, Boil-off, Giant Negative Magnetoresistance A. K. Ramadas and R. Rodriquez, Raman Scattering in Diluted Magnetic Semiconductors P. A. Wolff, Theory of Bound Magnetic Polarons in Semimagnetic Semiconductors

Volume 26 III–V Compound Semiconductors and Semiconductor Properties of Superionic Materials

Z. Yuanxi, III–V Compounds H. V. Winston, A. T. Hunter, H. Kimura, and R. E. Lee, InAs-Alloyed GaAs Substrates for Direct Implantation P. K. Bhattacharya and S. Dhar, Deep Levels in III–V Compound Semiconductors Grown by MBE Y. Ya. Gurevich and A. K. Ivanov-Shits, Semiconductor Properties of Supersonic Materials

Volume 27 High Conducting Quasi-One-Dimensional Organic Crystals

E. M. Conwell, Introduction to Highly Conducting Quasi-One-Dimensional Organic Crystals I. A. Howard, A Reference Guide to the Conducting Quasi-One-Dimensional Organic Molecular Crystals J. P. Pouqnet, Structural Instabilities E. M. Conwell, Transport Properties C. S. Jacobsen, Optical Properties J. C. Scolt, Magnetic Properties L. Zuppiroli, Irradiation Effects: Perfect Crystals and Real Crystals

Volume 28 Measurement of High-Speed Signals in Solid State Devices

J. Frey and D. Ioannou, Materials and Devices for High-Speed and Optoelectronic Applications H. Schumacher and E. Strid, Electronic Wafer Probing Techniques D. H. Auston, Picosecond Photoconductivity: High-Speed Measurements of Devices and Materials J. A. Valdmanis, Electro-Optic Measurement Techniques for Picosecond Materials, Devices and Integrated Circuits J. M. Wiesenfeld and R. K. Jain, Direct Optical Probing of Integrated Circuits and High-Speed Devices G. Plows, Electron-Beam Probing A. M. Weiner and R. B. Marcus, Photoemissive Probing 150 Contents of Volumes in this Series

Volume 29 Very High Speed Integrated Circuits: Gallium Arsenide LSI

M. Kuzuhara and T. Nazaki, Active Layer Formation by Ion Implantation H. Hasimoto, Focused Ion Beam Implantation Technology T. Nozaki and A. Higashisaka, Device Fabrication Process Technology M. Ino and T. Takada, GaAs LSI Circuit Design M. Hirayama, M. Ohmori, and K. Yamasaki, GaAs LSI Fabrication and Performance

Volume 30 Very High Speed Integrated Circuits: Heterostructure

H. Watanabe, T. Mizutani, and A. Usui, Fundamentals of Epitaxial Growth and Atomic Layer Epitaxy S. Hiyamizu, Characteristics of Two-Dimensional Electron Gas in III–V Compound Heterostructures Grown by MBE T. Nakanisi, Metalorganic Vapor Phase Epitaxy for High-Quality Active Layers T. Nimura, High Electron Mobility Transistor and LSI Applications T. Sugeta and T. Ishibashi, Hetero-Bipolar Transistor and LSI Application H. Matsuedo, T. Tanaka, and M. Nakamura, Optoelectronic Integrated Circuits

Volume 31 Indium Phosphide: Crystal Growth and Characterization

J. P. Farges, Growth of Discoloration-Free InP M. J. McCollum and G. E. Stillman, High Purity InP Grown by Hydride Vapor Phase Epitaxy I. Inada and T. Fukuda, Direct Synthesis and Growth of Indium Phosphide by the Liquid Phosphorous Encapsulated Czochralski Method O. Oda, K. Katagiri, K. Shinohara, S. Katsura, Y. Takahashi, K. Kainosho, K. Kohiro, and R. Hirano, InP Crystal Growth, Substrate Preparation and Evaluation K. Tada, M. Tatsumi, M. Morioka, T. Araki, and T. Kawase, InP Substrates: Production and Quality Control M. Razeghi, LP-MOCVD Growth, Characterization, and Application of InP Material T. A. Kennedy and P. J. Lin-Chung, Stoichiometric Defects in InP

Volume 32 Strained-Layer Superlattices: Physics

T. P. Pearsall, Strained-Layer Superlattices F. H. Pollack, Effects of Homogeneous Strain on the Electronic and Vibrational Levels in Semiconductors J. Y. Marzin, J. M. Gera´rd, P. Voisin, and J. A. Brum, Optical Studies of Strained III–V Heterolayers R. People and S. A. Jackson, Structurally Induced States from Strain and Confinement M. Jaros, Microscopic Phenomena in Ordered Superlattices

Volume 33 Strained-Layer Superlattices: Material Science and Technology

R. Hull and J. C. Bean, Principles and Concepts of Strained-Layer Epitaxy Contents of Volumes in this Series 151

W. J. Shaff, P. J. Tasker, M. C. Foisy, and L. F. Eastman, Device Applications of Strained-Layer Epitaxy S. T. Picraux, B. L. Doyle, and J. Y. Tsao, Structure and Characterization of Strained-Layer Superlattices E. Kasper and F. Schaffer, Group IV Compounds D. L. Martin, Molecular Beam Epitaxy of IV–VI Compounds Heterojunction R. L. Gunshor, L. A. Kolodziejski, A. V. Nurmikko, and N. Otsuka, Molecular Beam Epitaxy of I–VI Semiconductor Microstructures

Volume 34 Hydrogen in Semiconductors

J. I. Pankove and N. M. Johnson, Introduction to Hydrogen in Semiconductors C. H. Seager, Hydrogenation Methods J. I. Pankove, Hydrogenation of Defects in Crystalline Silicon J. W. Corbett, P. De´ak, U. V. Desnica, and S. J. Pearton, Hydrogen Passivation of Damage Centers in Semiconductors S. J. Pearton, Neutralization of Deep Levels in Silicon J. I. Pankove, Neutralization of Shallow Acceptors in Silicon N. M. Johnson, Neutralization of Donor Dopants and Formation of Hydrogen-Induced Defects in n-Type Silicon M. Stavola and S. J. Pearton, Vibrational Spectroscopy of Hydrogen-Related Defects in Silicon A. D. Marwick, Hydrogen in Semiconductors: Ion Beam Techniques C. Herring and N. M. Johnson, Hydrogen Migration and Solubility in Silicon E. E. Haller, Hydrogen-Related Phenomena in Crystalline Germanium J. Kakalios, Hydrogen Diffusion in Amorphous Silicon J. Chevalier, B. Clerjaud, and B. Pajot, Neutralization of Defects and Dopants in III–V Semiconductors G. G. DeLeo and W. B. Fowler, Computational Studies of Hydrogen-Containing Complexes in Semiconductors R. F. Kiefl and T. L. Estle, Muonium in Semiconductors C. G. Van de Walle, Theory of Isolated Interstitial Hydrogen and Muonium in Crystalline Semiconductors

Volume 35 Nanostructured Systems

M. Reed, Introduction H. van Houten, C. W. J. Beenakker, and B. J. Wees, Quantum Point Contacts G. Timp, When Does a Wire Become an Electron Waveguide? M. Bu´ttiker, The Quantum Hall Effects in Open Conductors W. Hansen, J. P. Kotthaus, and U. Merkt, Electrons in Laterally Periodic Nanostructures

Volume 36 The Spectroscopy of Semiconductors

D. Heiman, Spectroscopy of Semiconductors at Low Temperatures and High Magnetic Fields A. V. Nurmikko, Transient Spectroscopy by Ultrashort Laser Pulse Techniques 152 Contents of Volumes in this Series

A. K. Ramdas and S. Rodriguez, Piezospectroscopy of Semiconductors O. J. Glembocki and B. V. Shanabrook, Photoreflectance Spectroscopy of Microstructures D. G. Seiler, C. L. Littler, and M. H. Wiler, One- and Two-Photon Magneto-Optical Spectroscopy of

InSb and Hg1ÀxCdx Te

Volume 37 The Mechanical Properties of Semiconductors

A.-B. Chen, A. Sher, and W. T. Yost, Elastic Constants and Related Properties of Semiconductor Compounds and Their Alloys D. R. Clarke, Fracture of Silicon and Other Semiconductors H. Siethoff, The Plasticity of Elemental and Compound Semiconductors S. Guruswamy, K. T. Faber, and J. P. Hirth, Mechanical Behavior of Compound Semiconductors S. Mahajan, Deformation Behavior of Compound Semiconductors J. P. Hirth, Injection of Dislocations into Strained Multilayer Structures D. Kendall, C. B. Fleddermann, and K. J. Malloy, Critical Technologies for the Micromatching of Silicon J. Matsuba and K. Mokuya, Processing and Semiconductor Thermoelastic Behavior

Volume 38 Imperfections in III/V Materials

U. Scherz and M. Scheffler, Density-Functional Theory of sp-Bonded Defects in III/V Semiconductors M. Kaminska and E. R. Weber, E12 Defect in GaAs D. C. Look, Defects Relevant for Compensation in Semi-Insulating GaAs R. C. Newman, Local Vibrational Mode Spectroscopy of Defects in III/V Compounds A. M. Hennel, Transition Metals in III/V Compounds K. J. Malloy and K. Khachaturyan, DX and Related Defects in Semiconductors V. Swaminathan and A. S. Jordan, Dislocations in III/V Compounds K. W. Nauka, Deep Level Defects in the Epitaxial III/V Materials

Volume 39 Minority Carriers in III–V Semiconductors: Physics and Applications

N. K. Dutta, Radiative Transition in GaAs and Other III–V Compounds R. K. Ahrenkiel, Minority-Carrier Lifetime in III–V Semiconductors T. Furuta, High Field Minority Electron Transport in p-GaAs M. S. Lundstrom, Minority-Carrier Transport in III–V Semiconductors R. A. Abram, Effects of Heavy Doping and High Excitation on the Band Structure of GaAs D. Yevick and W. Bardyszewski, An Introduction to Non-Equilibrium Many-Body Analyses of Optical Processes in III–V Semiconductors

Volume 40 Epitaxial Microstructures

E. F. Schubert, Delta-Doping of Semiconductors: Electronic, Optical and Structural Properties of Materials and Devices A. Gossard, M. Sundaram, and P. Hopkins, Wide Graded Potential Wells Contents of Volumes in this Series 153

P. Petroff, Direct Growth of Nanometer-Size Quantum Wire Superlattices E. Kapon, Lateral Patterning of Quantum Well Heterostructures by Growth of Nonplanar Substrates

H. Temkin, D. Gershoni, and M. Panish, Optical Properties of Ga1ÀxInxAs/InP Quantum Wells

Volume 41 High Speed Heterostructure Devices

F. Capasso, F. Beltram, S. Sen, A. Pahlevi, and A. Y. Cho, Quantum Electron Devices: Physics and Applications P. Solomon, D. J. Frank, S. L. Wright and F. Canora, GaAs-Gate Semiconductor-Insulator- Semiconductor FET M. H. Hashemi and U. K. Mishra, Unipolar InP-Based Transistors R. Kiehl, Complementary Heterostructure FET Integrated Circuits T. Ishibashi, GaAs-Based and InP-Based Heterostructure Bipolar-Transistors H. C. Liu and T. C. L. G. Sollner, High-Frequency-Tunneling Devices H. Ohnishi, T. More, M. Takatsu, K. Imamura, and N. Yokoyama, Resonant-Tunneling Hot-Electron Transistors and Circuits

Volume 42 Oxygen in Silicon

F. Shimura, Introduction to Oxygen in Silicon W. Lin, The Incorporation of Oxygen into Silicon Crystals T. J. Schaffner and D. K. Schroder, Characterization Techniques for Oxygen in Silicon W. M. Bullis, Oxygen Concentration Measurement S. M. Hu, Intrinsic Point Defects in Silicon B. Pajot, Some Atomic Configuration of Oxygen J. Michel and L. C. Kimerling, Electrical Properties of Oxygen in Silicon R. C. Newman and R. Jones, Diffusion of Oxygen in Silicon T. Y. Tan and W. J. Taylor, Mechanisms of Oxygen Precipitation: Some Quantitative Aspects M. Schrems, Simulation of Oxygen Precipitation K. Simino and I. Yonenaga, Oxygen Effect on Mechanical Properties W. Bergholz, Grown-in and Process-Induced Effects F. Shimura, Intrinsic/Internal Gettering H. Tsuya, Oxygen Effect on Electronic Device Performance

Volume 43 Semiconductors for Room Temperature Nuclear Detector Applications

R. B. James and T. E. Schlesinger, Introduction and Overview L. S. Darken and C. E. Cox, High-Purity Germanium Detectors A. Burger, D. Nason, L. Van den Berg, and M. Schieber, Growth of Mercuric Iodide X. J. Bao, T. E. Schlesinger, and R. B. James, Electrical Properties of Mercuric Iodide X. J. Bao, R. B. James, and T. E. Schlesinger, Optical Properties of Red Mercuric Iodide M. Hage-Ali and P. Siffert, Growth Methods of CdTe Nuclear Detector Materials M. Hage-Ali and P. Siffert, Characterization of CdTe Nuclear Detector Materials 154 Contents of Volumes in this Series

M. Hage-Ali and P. Siffert, CdTe Nuclear Detectors and Applications

R. B. James, T. E. Schlesinger, J. Lund, and M. Schieber,Cd1Àx Znx Te Spectrometers for Gamma and X-Ray Applications D. S. McGregor, J. E. Kammeraad, Gallium Arsenide Radiation Detectors and Spectrometers J. C. Lund, F. Olschner, and A. Burger, Lead Iodide M. R. Squillante and K. S. Shah, Other Materials: Status and Prospects V. M. Gerrish, Characterization and Quantification of Detector Performance J. S. Iwanczyk and B. E. Patt, Electronics for X-ray and Gamma Ray Spectrometers M. Schieber, R. B. James and T. E. Schlesinger, Summary and Remaining Issues for Room Temperature Radiation Spectrometers

Volume 44 II–IV Blue/Green Light Emitters: Device Physics and Epitaxial Growth

J. Han and R. L. Gunshor, MBE Growth and Electrical Properties of Wide Bandgap ZnSe-based II–VI Semiconductors S. Fujita and S. Fujita, Growth and Characterization of ZnSe-based II–VI Semiconductors by MOVPE E. Ho and L. A. Kolodziejski, Gaseous Source UHV Epitaxy Technologies for Wide Bandgap II–VI Semiconductors C. G. Van de Walle, Doping of Wide-Band-Gap II–VI Compounds – Theory R. Cingolani, Optical Properties of Excitons in ZnSe-Based Quantum Well Heterostructures A. Ishibashi and A. V. Nurmikko, II–VI Diode Lasers: A Current View of Device Performance and Issues S. Guha and J. Petruzello, Defects and Degradation in Wide-Gap II–VI-based Structure and Light Emitting Devices

Volume 45 Effect of Disorder and Defects in Ion-Implanted Semiconductors: Electrical and Physiochemical Characterization

H. Ryssel, Ion Implantation into Semiconductors: Historical Perspectives You-Nian Wang and Teng-Cai Ma, Electronic Stopping Power for Energetic Ions in Solids S. T. Nakagawa, Solid Effect on the Electronic Stopping of Crystalline Target and Application to Range Estimation G. Miller, S. Kalbitzer, and G. N. Greaves, Ion Beams in Amorphous Semiconductor Research J. Boussey-Said, Sheet and Spreading Resistance Analysis of Ion Implanted and Annealed Semiconductors M. L. Polignano and G. Queirolo, Studies of the Stripping Hall Effect in Ion-Implanted Silicon J. Sroemenos, Transmission Electron Microscopy Analyses R. Nipoti and M. Servidori, Rutherford Backscattering Studies of Ion Implanted Semiconductors P. Zaumseil, X-ray Diffraction Techniques

Volume 46 Effect of Disorder and Defects in Ion-Implanted Semiconductors: Optical and Photothermal Characterization

M. Fried, T. Lohner, and J. Gyulai, Ellipsometric Analysis A. Seas and C. Christofides, Transmission and Reflection Spectroscopy on Ion Implanted Semiconductors Contents of Volumes in this Series 155

A. Othonos and C. Christofides, Photoluminescence and Raman Scattering of Ion Implanted Semiconductors. Influence of Annealing C. Christofides, Photomodulated Thermoreflectance Investigation of Implanted Wafers. Annealing Kinetics of Defects U. Zammit, Photothermal Deflection Spectroscopy Characterization of Ion-Implanted and Annealed Silicon Films A. Mandelis, A. Budiman, and M. Vargas, Photothermal Deep-Level Transient Spectroscopy of Impurities and Defects in Semiconductors R. Kalish and S. Charbonneau, Ion Implantation into Quantum-Well Structures A. M. Myasnikov and N. N. Gerasimenko, Ion Implantation and Thermal Annealing of III–V Compound Semiconducting Systems: Some Problems of III–V Narrow Gap Semiconductors

Volume 47 Uncooled Infrared Imaging Arrays and Systems

R. G. Buser and M. P. Tompsett, Historical Overview P. W. Kruse, Principles of Uncooled Infrared Focal Plane Arrays R. A. Wood, Monolithic Silicon Microbolometer Arrays C. M. Hanson, Hybrid Pyroelectric-Ferroelectric Bolometer Arrays D. L. Polla and J. R. Choi, Monolithic Pyroelectric Bolometer Arrays N. Teranishi, Thermoelectric Uncooled Infrared Focal Plane Arrays M. F. Tompsett, Pyroelectric Vidicon T. W. Kenny, Tunneling Infrared Sensors J. R. Vig, R. L Filler, and Y. Kim, Application of Quartz Microresonators to Uncooled Infrared Imaging Arrays P. W. Kruse, Application of Uncooled Monolithic Thermoelectric Linear Arrays to Imaging Radiometers

Volume 48 High Brightness Light Emitting Diodes

G. B. Stringfellow, Materials Issues in High-Brightness Light-Emitting Diodes M. G. Craford, Overview of Device Issues in High-Brightness Light-Emitting Diodes F. M. Steranka, AlGaAs Red Light Emitting Diodes C. H. Chen, S. A. Stockman, M. J. Peanasky, and C. P. Kuo, OMVPE Growth of AlGaInP for High Efficiency Visible Light-Emitting Diodes F. A. Kish and R. M. Fletcher, AlGaInP Light-Emitting Diodes M. W. Hodapp, Applications for High Brightness Light-Emitting Diodes J. Akasaki and H. Amano, Organometallic Vapor Epitaxy of GaN for High Brightness Blue Light Emitting Diodes S. Nakamura, Group III–V Nitride Based Ultraviolet-Blue-Green-Yellow Light-Emitting Diodes and Laser Diodes

Volume 49 Light Emission in Silicon: from Physics to Devices

D. J. Lockwood, Light Emission in Silicon G. Abstreiter, Band Gaps and Light Emission in Si/SiGe Atomic Layer Structures 156 Contents of Volumes in this Series

T. G. Brown and D. G. Hall, Radiative Isoelectronic Impurities in Silicon and Silicon-Germanium Alloys and Superlattices J. Michel, L. V. C. Assali, M. T. Morse, and L. C. Kimerling, Erbium in Silicon Y. Kanemitsu, Silicon and Germanium Nanoparticles P. M. Fauchet, Porous Silicon: Photoluminescence and Electroluminescent Devices C. Delerue, G. Allan, and M. Lannoo, Theory of Radiative and Nonradiative Processes in Silicon Nanocrystallites L. Brus, Silicon Polymers and Nanocrystals

Volume 50 Gallium Nitride (GaN)

J. I. Pankove and T. D. Moustakas, Introduction S. P. DenBaars and S. Keller, Metalorganic Chemical Vapor Deposition (MOCVD) of Group III Nitrides W. A. Bryden and T. J. Kistenmacher, Growth of Group III–A Nitrides by Reactive Sputtering N. Newman, Thermochemistry of III–N Semiconductors S. J. Pearton and R. J. Shul, Etching of III Nitrides S. M. Bedair, Indium-based Nitride Compounds A. Trampert, O. Brandt, and K. H. Ploog, Crystal Structure of Group III Nitrides H. Morkoc¸, F. Hamdani, and A. Salvador, Electronic and Optical Properties of III–V Nitride based Quantum Wells and Superlattices K. Doverspike and J. I. Pankove, Doping in the III-Nitrides T. Suski and P. Perlin, High Pressure Studies of Defects and Impurities in Gallium Nitride B. Monemar, Optical Properties of GaN W. R. L. Lambrecht, Band Structure of the Group III Nitrides N. E. Christensen and P. Perlin, Phonons and Phase Transitions in GaN S. Nakamura, Applications of LEDs and LDs I. Akasaki and H. Amano, Lasers J. A. Cooper, Jr., Nonvolatile Random Access Memories in Wide Bandgap Semiconductors

Volume 51A Identification of Defects in Semiconductors

G. D. Watkins, EPR and ENDOR Studies of Defects in Semiconductors J.-M. Spaeth, Magneto-Optical and Electrical Detection of Paramagnetic Resonance in Semiconductors T. A. Kennedy and E. R. Claser, Magnetic Resonance of Epitaxial Layers Detected by Photoluminescence K. H. Chow, B. Hitti, and R. F. Kiefl, μSR on Muonium in Semiconductors and Its Relation to Hydrogen K. Saarinen, P. Hautoja¨rvi, and C. Corbel, Positron Annihilation Spectroscopy of Defects in Semiconductors R. Jones and P. R. Briddon, The Ab Initio Cluster Method and the Dynamics of Defects in Semiconductors

Volume 51B Identification Defects in Semiconductors

G. Davies, Optical Measurements of Point Defects P. M. Mooney, Defect Identification Using Capacitance Spectroscopy Contents of Volumes in this Series 157

M. Stavola, Vibrational Spectroscopy of Light Element Impurities in Semiconductors P. Schwander, W. D. Rau, C. Kisielowski, M. Gribelyuk, and A. Ourmazd, Defect Processes in Semiconductors Studied at the Atomic Level by Transmission Electron Microscopy N. D. Jager and E. R. Weber, Scanning Tunneling Microscopy of Defects in Semiconductors

Volume 52 SiC Materials and Devices

K. Ja¨rrendahl and R. F. Davis, Materials Properties and Characterization of SiC V. A. Dmitiriev and M. G. Spencer, SiC Fabrication Technology: Growth and Doping V. Saxena and A. J. Steckl, Building Blocks for SiC Devices: Ohmic Contacts, Schottky Contacts, and p-n Junctions M. S. Shur, SiC Transistors C. D. Brandt, R. C. Clarke, R. R. Siergiej, J. B. Casady, A. W. Morse, S. Sriram, and A. K. Agarwal, SiC for Applications in High-Power Electronics R. J. Trew, SiC Microwave Devices J. Edmond, H. Kong, G. Negley, M. Leonard, K. Doverspike, W. Weeks, A. Suvorov, D. Waltz, and C. Carter, Jr., SiC-Based UV Photodiodes and Light-Emitting Diodes H. Morkoc¸, Beyond Silicon Carbide! III–V Nitride-Based Heterostructures and Devices

Volume 53 Cumulative Subjects and Author Index Including Tables of Contents for Volumes 1–50

Volume 54 High Pressure in Semiconductor Physics I

W. Paul, High Pressure in Semiconductor Physics: A Historical Overview N. E. Christensen, Electronic Structure Calculations for Semiconductors Under Pressure R. J. Neimes and M. I. McMahon, Structural Transitions in the Group IV, III–V and II–VI Semiconductors Under Pressure A. R. Goni and K. Syassen, Optical Properties of Semiconductors Under Pressure P. Trautman, M. Baj, and J. M. Baranowski, Hydrostatic Pressure and Uniaxial Stress in Investigations of the EL2 Defect in GaAs M. Li and P. Y. Yu, High-Pressure Study of DX Centers Using Capacitance Techniques T. Suski, Spatial Correlations of Impurity Charges in Doped Semiconductors N. Kuroda, Pressure Effects on the Electronic Properties of Diluted Magnetic Semiconductors

Volume 55 High Pressure in Semiconductor Physics II

D. K. Maude and J. C. Portal, Parallel Transport in Low-Dimensional Semiconductor Structures P. C. Klipstein, Tunneling Under Pressure: High-Pressure Studies of Vertical Transport in Semiconductor Heterostructures E. Anastassakis and M. Cardona, Phonons, Strains, and Pressure in Semiconductors 158 Contents of Volumes in this Series

F. H. Pollak, Effects of External Uniaxial Stress on the Optical Properties of Semiconductors and Semiconductor Microstructures A. R. Adams, M. Silver, and J. Allam, Semiconductor Optoelectronic Devices S. Porowski and I. Grzegory, The Application of High Nitrogen Pressure in the Physics and Technology of III–N Compounds M. Yousuf, Diamond Anvil Cells in High Pressure Studies of Semiconductors

Volume 56 Germanium Silicon: Physics and Materials

J. C. Bean, Growth Techniques and Procedures D. E. Savage, F. Liu, V. Zielasek, and M. G. Lagally, Fundamental Crystal Growth Mechanisms R. Hull, Misfit Strain Accommodation in SiGe Heterostructures M. J. Shaw and M. Jaros, Fundamental Physics of Strained Layer GeSi: Quo Vadis? F. Cerdeira, Optical Properties S. A. Ringel and P. N. Grillot, Electronic Properties and Deep Levels in Germanium-Silicon J. C. Campbell, Optoelectronics in Silicon and Germanium Silicon

K. Eberl, K. Brunner, and O. G. Schmidt,Si1ÀyCy and Si1ÀxÀyGe2Cy Alloy Layers

Volume 57 Gallium Nitride (GaN) II

R. J. Molnar, Hydride Vapor Phase Epitaxial Growth of III–V Nitrides T. D. Moustakas, Growth of III–V Nitrides by Molecular Beam Epitaxy Z. Liliental-Weber, Defects in Bulk GaN and Homoepitaxial Layers C. G. Van de Walk and N. M. Johnson, Hydrogen in III–V Nitrides W. Götz and N. M. Johnson, Characterization of Dopants and Deep Level Defects in Gallium Nitride B. Gil, Stress Effects on Optical Properties C. Kisielowski, Strain in GaN Thin Films and Heterostructures J. A. Miragliotta and D. K. Wickenden, Nonlinear Optical Properties of Gallium Nitride B. K. Meyer, Magnetic Resonance Investigations on Group III–Nitrides M. S. Shur and M. Asif Khan, GaN and AIGaN Ultraviolet Detectors C. H. Qiu, J. I. Pankove, and C. Rossington, II–V Nitride-Based X-ray Detectors

Volume 58 Nonlinear Optics in Semiconductors I

A. Kost, Resonant Optical Nonlinearities in Semiconductors E. Garmire, Optical Nonlinearities in Semiconductors Enhanced by Carrier Transport D. S. Chemla, Ultrafast Transient Nonlinear Optical Processes in Semiconductors M. Sheik-Bahae and E. W. Van Stryland, Optical Nonlinearities in the Transparency Region of Bulk Semiconductors J. E. Millerd, M. Ziari, and A. Partovi, Photorefractivity in Semiconductors Contents of Volumes in this Series 159

Volume 59 Nonlinear Optics in Semiconductors II

J. B. Khurgin, Second Order Nonlinearities and Optical Rectification K. L. Hall, E. R. Thoen, and E. P. Ippen, Nonlinearities in Active Media E. Hanamura, Optical Responses of Quantum Wires/Dots and Microcavities U. Keller, Semiconductor Nonlinearities for Solid-State Laser Modelocking and Q-Switching A. Miller, Transient Grating Studies of Carrier Diffusion and Mobility in Semiconductors

Volume 60 Self-Assembled InGaAs/GaAs Quantum Dots

Mitsuru Sugawara, Theoretical Bases of the Optical Properties of Semiconductor Quantum Nano- Structures Yoshiaki Nakata, Yoshihiro Sugiyama, and Mitsuru Sugawara, Molecular Beam Epitaxial Growth of Self- Assembled InAs/GaAs Quantum Dots Kohki Mukai, Mitsuru Sugawara, Mitsuru Egawa, and Nobuyuki Ohtsuka, Metalorganic Vapor Phase Epitaxial Growth of Self-Assembled InGaAs/GaAs Quantum Dots Emitting at 1.3 μm Kohki Mukai and Mitsuru Sugawara, Optical Characterization of Quantum Dots Kohki Mukai and Milsuru Sugawara, The Photon Bottleneck Effect in Quantum Dots Hajime Shoji, Self-Assembled Quantum Dot Lasers Hiroshi Ishikawa, Applications of Quantum Dot to Optical Devices Mitsuru Sugawara, Kohki Mukai, Hiroshi Ishikawa, Koji Otsubo, and Yoshiaki Nakata, The Latest News

Volume 61 Hydrogen in Semiconductors II

Norbert H. Nickel, Introduction to Hydrogen in Semiconductors II Noble M. Johnson and Chris G. Van de Walle, Isolated Monatomic Hydrogen in Silicon Yurij V. Gorelkinskii, Electron Paramagnetic Resonance Studies of Hydrogen and Hydrogen-Related Defects in Crystalline Silicon Norbert H. Nickel, Hydrogen in Polycrystalline Silicon Wolfhard Beyer, Hydrogen Phenomena in Hydrogenated Amorphous Silicon Chris G. Van de Walle, Hydrogen Interactions with Polycrystalline and Amorphous Silicon–Theory Karen M. McManus Rutledge, Hydrogen in Polycrystalline CVD Diamond Roger L. Lichti, Dynamics of Muonium Diffusion, Site Changes and Charge-State Transitions Matthew D. McCluskey and Eugene E. Haller, Hydrogen in III–V and II–VI Semiconductors S. J. Pearton and J. W. Lee, The Properties of Hydrogen in GaN and Related Alloys Jörg Neugebauer and Chris G. Van de Walle, Theory of Hydrogen in GaN

Volume 62 Intersubband Transitions in Quantum Wells: Physics and Device Applications I

Manfred Helm, The Basic Physics of Intersubband Transitions Jerome Faist, Carlo Sirtori, Federico Capasso, Loren N. Pfeiffer, Ken W. West, Deborah L. Sivco, and Alfred Y. Cho, Quantum Interference Effects in Intersubband Transitions H. C. Liu, Quantum Well Infrared Photodetector Physics and Novel Devices S. D. Gunapala and S. V. Bandara, Quantum Well Infrared Photodetector (QWIP) Focal Plane Arrays 160 Contents of Volumes in this Series

Volume 63 Chemical Mechanical Polishing in Si Processing

Frank B. Kaufman, Introduction Thomas Bibby and Karey Holland, Equipment John P. Bare, Facilitization Duane S. Boning and Okumu Ouma, Modeling and Simulation Shin Hwa Li, Bruce Tredinnick, and Mel Hoffman, Consumables I: Slurry Lee M. Cook, CMP Consumables II: Pad Franc¸ois Tardif, Post-CMP Clean Shin Hwa Li, Tara Chhatpar, and Frederic Robert, CMP Metrology Shin Hwa Li, Visun Bucha, and Kyle Wooldridge, Applications and CMP-Related Process Problems

Volume 64 Electroluminescence I

M. G. Craford, S. A. Stockman, M. J. Peansky, and F. A. Kish, Visible Light-Emitting Diodes H. Chui, N. F. Gardner, P. N. Grillot, J. W. Huang, M. R. Krames, and S. A. Maranowski, High-Efficiency AIGaInP Light-Emitting Diodes R. S. Kern, W. Go¯tz, C. H. Chen, H. Liu, R. M. Fletcher, and C. P. Kuo, High-Brightness Nitride-Based Visible-Light-Emitting Diodes Yoshiharu Sato, Organic LED System Considerations V. Bulovic´, P. E. Burrows, and S. R. Forrest, Molecular Organic Light-Emitting Devices

Volume 65 Electroluminescence II

V. Bulovic´ and S. R. Forrest, Polymeric and Molecular Organic Light Emitting Devices: A Comparison Regina Mueller-Mach and Gerd O. Mueller, Thin Film Electroluminescence Markku Leskela¯, Wei-Min Li, and Mikko Ritala, Materials in Thin Film Electroluminescent Devices Kristiaan Neyts, Microcavities for Electroluminescent Devices

Volume 66 Intersubband Transitions in Quantum Wells: Physics and Device Applications II

Jerome Faist, Federico Capasso, Carlo Sirtori, Deborah L. Sivco, and Alfred Y. Cho, Quantum Cascade Lasers Federico Capasso, Carlo Sirtori, D. L. Sivco, and A. Y. Cho, Nonlinear Optics in Coupled-Quantum- Well Quasi-Molecules Karl Unterrainer, Photon-Assisted Tunneling in Semiconductor Quantum Structures P. Haring Bolivar, T. Dekorsy, and H. Kurz, Optically Excited Bloch Oscillations–Fundamentals and Application Perspectives

Volume 67 Ultrafast Physical Processes in Semiconductors

Alfred Leitenstorfer and Alfred Laubereau, Ultrafast Electron-Phonon Interactions in Semiconductors: Quantum Kinetic Memory Effects Contents of Volumes in this Series 161

Christoph Lienau and Thomas Elsaesser, Spatially and Temporally Resolved Near-Field Scanning Optical Microscopy Studies of Semiconductor Quantum Wires K. T. Tsen, Ultrafast Dynamics in Wide Bandgap Wurtzite GaN J. Paul Callan, Albert M.-T. Kim, Christopher A. D. Roeser, and Eriz Mazur, Ultrafast Dynamics and Phase Changes in Highly Excited GaAs Hartmut Hang, Quantum Kinetics for Femtosecond Spectroscopy in Semiconductors T. Meier and S. W. Koch, Coulomb Correlation Signatures in the Excitonic Optical Nonlinearities of Semiconductors Roland E. Allen, Traian Dumitrica˘, and Ben Torralva, Electronic and Structural Response of Materials to Fast, Intense Laser Pulses E. Gornik and R. Kersting, Coherent THz Emission in Semiconductors

Volume 68 Isotope Effects in Solid State Physics

Vladimir G. Plekhanov, Elastic Properties; Thermal Properties; Vibrational Properties; Raman Spectra of Isotopically Mixed Crystals; Excitons in LiH Crystals; Exciton–Phonon Interaction; Isotopic Effect in the Emission Spectrum of Polaritons; Isotopic Disordering of Crystal Lattices; Future Developments and Applications; Conclusions

Volume 69 Recent Trends in Thermoelectric Materials Research I

H. Julian Goldsmid, Introduction Terry M. Tritt and Valerie M. Browning, Overview of Measurement and Characterization Techniques for Thermoelectric Materials Mercouri G. Kanatzidis, The Role of Solid-State Chemistry in the Discovery of New Thermoelectric Materials B. Lenoir, H. Scherrer, and T. Caillat, An Overview of Recent Developments for BiSb Alloys Citrad Uher, Skutterudities: Prospective Novel Thermoelectrics George S. Nolas, Glen A. Slack, and Sandra B. Schujman, Semiconductor Clathrates: A Phonon Glass Electron Crystal Material with Potential for Thermoelectric Applications

Volume 70 Recent Trends in Thermoelectric Materials Research II

Brian C. Sales, David G. Mandrus, and Bryan C. Chakoumakos, Use of Atomic Displacement Parameters in Thermoelectric Materials Research S. Joseph Poon, Electronic and Thermoelectric Properties of Half-Heusler Alloys Terry M. Tritt, A. L. Pope, and J. W. Kolis, Overview of the Thermoelectric Properties of Quasicrystalline Materials and Their Potential for Thermoelectric Applications Alexander C. Ehrlich and Stuart A. Wolf, Military Applications of Enhanced Thermoelectrics David J. Singh, Theoretical and Computational Approaches for Identifying and Optimizing Novel Thermoelectric Materials Terry M. Tritt and R. T. Littleton, IV, Thermoelectric Properties of the Transition Metal Pentatellurides: Potential Low-Temperature Thermoelectric Materials 162 Contents of Volumes in this Series

Franz Freibert, Timothy W. Darling, Albert Miglori, and Stuart A. Trugman, Thermomagnetic Effects and Measurements M. Bartkowiak and G. D. Mahan, Heat and Electricity Transport Through Interfaces

Volume 71 Recent Trends in Thermoelectric Materials Research III

M. S. Dresselhaus, Y.-M. Lin, T. Koga, S. B. Cronin, O. Rabin, M. R. Black, and G. Dresselhaus, Quantum Wells and Quantum Wires for Potential Thermoelectric Applications D. A. Broido and T. L. Reinecke, Thermoelectric Transport in Quantum Well and Quantum Wire Superlattices G. D. Mahan, Thermionic Refrigeration Rama Venkatasubramanian, Phonon Blocking Electron Transmitting Superlattice Structures as Advanced Thin Film Thermoelectric Materials G. Chen, Phonon Transport in Low-Dimensional Structures

Volume 72 Silicon Epitaxy

S. Acerboni, ST Microelectronics, CFM-AGI Department, Agrate Brianza, Italy V.-M. Airaksinen, Okmetic Oyj R&D Department, Vantaa, Finland G. Beretta, ST Microelectronics, DSG Epitaxy Catania Department, Catania, Italy C. Cavallotti, Dipartimento di Chimica Fisica Applicata, Politecnico di Milano, Milano, Italy D. Crippa, MEMC Electronic Materials, Epitaxial and CVD Department, Operations Technology Division, Novara, Italy D. Dutartre, ST Microelectronics, Central R&D, Crolles, France Srikanth Kommu, MEMC Electronic Materials inc., EPI Technology Group, St. Peters, Missouri M. Masi, Dipartimento di Chimica Fisica Applicata, Politecnico di Milano, Milano, Italy D. J. Meyer, ASM Epitaxy, Phoenix, Arizona J. Murota, Research Institute of Electrical Communication, Laboratory for Electronic Intelligent Systems, Tohoku University, Sendai, Japan V. Pozzetti, LPE Epitaxial Technologies, Bollate, Italy A. M. Rinaldi, MEMC Electronic Materials, Epitaxial and CVD Department, Operations Technology Division, Novara, Italy Y. Shiraki, Research Center for Advanced Science and Technology (RCAST), University of Tokyo, Tokyo, Japan

Volume 73 Processing and Properties of Compound Semiconductors

S. J. Pearton, Introduction Eric Donkor, Gallium Arsenide Heterostructures Annamraju Kasi Viswanatli, Growth and Optical Properties of GaN D. Y. C. Lie and K. L. Wang, SiGe/Si Processing S. Kim and M. Razeghi, Advances in Quantum Dot Structures Walter P. Gomes, Wet Etching of III–V Semiconductors Contents of Volumes in this Series 163

Volume 74 Silicon-Germanium Strained Layers and Heterostructures

S. C. Jain and M. Willander, Introduction; Strain, Stability, Reliability and Growth; Mechanism of Strain Relaxation; Strain, Growth, and TED in SiGeC Layers; Bandstructure and Related Properties; Heterostructure Bipolar Transistors; FETs and Other Devices

Volume 75 Laser Crystallization of Silicon

Norbert H. Nickel, Introduction to Laser Crystallization of Silicon Costas P. Grigoropoidos, Seung-Jae Moon and Ming-Hong Lee, Heat Transfer and Phase Transformations in Laser Melting and Recrystallization of Amorphous Thin Si Films ˇ Robert Cerny´ and Petr Prˇikryl, Modeling Laser-Induced Phase-Change Processes: Theory and Computation Paulo V. Santos, Laser Interference Crystallization of Amorphous Films Philipp Lengsfeld and Norbert H. Nickel, Structural and Electronic Properties of Laser-Crystallized Poly-Si

Volume 76 Thin-Film Diamond I

X. Jiang, Textured and Heteroepitaxial CVD Diamond Films Eberhard Blank, Structural Imperfections in CVD Diamond Films R. Kalish, Doping Diamond by Ion-Implantation A. Deneuville, Boron Doping of Diamond Films from the Gas Phase S. Koizumi, n-Type Diamond Growth C. E. Nebel, Transport and Defect Properties of Intrinsic and Boron-Doped Diamond Milosˇ Nesla´dek, Ken Haenen and Milan Vaneˇcˇek, Optical Properties of CVD Diamond Rolf Sauer, Luminescence from Optical Defects and Impurities in CVD Diamond

Volume 77 Thin-Film Diamond II

Jacques Chevallier, Hydrogen Diffusion and Acceptor Passivation in Diamond Jurgen€ Ristein, Structural and Electronic Properties of Diamond Surfaces John C. Angus, Yuri V. Pleskov and Sally C. Eaton, Electrochemistry of Diamond Greg M. Swain, Electroanalytical Applications of Diamond Electrodes Werner Haenni, Philippe Rychen, Matthyas Fryda and Christos Comninellis, Industrial Applications of Diamond Electrodes Philippe Bergonzo and Richard B. Jackman, Diamond-Based Radiation and Photon Detectors Hiroshi Kawarada, Diamond Field Effect Transistors Using H-Terminated Surfaces Shinichi Shikata and Hideaki Nakahata, Diamond Surface Acoustic Wave Device

Volume 78 Semiconducting Chalcogenide Glass I

V. S. Minaev and S. P. Timoshenkov, Glass-Formation in Chalcogenide Systems and Periodic System A. Popov, Atomic Structure and Structural Modification of Glass 164 Contents of Volumes in this Series

V. A. Funtikov, Eutectoidal Concept of Glass Structure and Its Application in Chalcogenide Semiconductor Glasses V. S. Minaev, Concept of Polymeric Polymorphous-Crystalloid Structure of Glass and Chalcogenide Systems: Structure and Relaxation of Liquid and Glass

Volume 79 Semiconducting Chalcogenide Glass II

M. D. Bal’makov, Information Capacity of Condensed Systems ˇ A. Cesnys, G. Jusˇka and E. Montrimas, Charge Carrier Transfer at High Electric Fields in Noncrystalline Semiconductors Andrey S. Glebov, The Nature of the Current Instability in Chalcogenide Vitreous Semiconductors A. M. Andriesh, M. S. Iovu and S. D. Shutov, Optical and Photoelectrical Properties of Chalcogenide Glasses V. Val. Sobolev and V. V. Sobolev, Optical Spectra of Arsenic Chalcogenides in a Wide Energy Range of Fundamental Absorption Yu. S. Tver’yanovich, Magnetic Properties of Chalcogenide Glasses

Volume 80 Semiconducting Chalcogenide Glass III

Andrey S. Glebov, Electronic Devices and Systems Based on Current Instability in Chalcogenide Semiconductors Dumitru Tsiulyanu, Heterostructures on Chalcogenide Glass and Their Applications E. Bychkov, Yu. Tveryanovich and Yu. Vlasov, Ion Conductivity and Sensors Yu. S. Tver’yanovich and A. Tverjanovich, Rare-earth Doped Chalcogenide Glass M. F. Churbanov and V. G. Plotnichenko, Optical Fibers from High-purity Arsenic Chalcogenide Glasses

Volume 81 Conducting Organic Materials and Devices

Suresh C. Jain, Magnus Willander and Vikram Kumar, Introduction; Polyacetylene; Optical and Transport Properties; Light Emitting Diodes and Lasers; Solar Cells; Transistors

Volume 82 Semiconductors and Semimetals

Maiken H. Mikkelsen, Roberto C. Myers, Gregory D. Fuchs, and David D. Awschalom, Single Spin Coherence in Semiconductors Jairo Sinova and A. H. MacDonald, Theory of Spin–Orbit Effects in Semiconductors K. M. Yu, T. Wojtowicz, W. Walukiewicz, X. Liu, and J. K. Furdyna, Fermi Level Effects on Mn Incorporation in III–Mn–V Ferromagnetic Semiconductors T. Jungwirth, B. L. Gallagher, and J.Wunderlich, Transport Properties of Ferromagnetic Semiconductors F. Matsukura, D. Chiba, and H. Ohno, Spintronic Properties of Ferromagnetic Semiconductors C. Gould, G. Schmidt, and L. W. Molenkamp, Spintronic Nanodevices Contents of Volumes in this Series 165

J. Cibert, L. Besombes, D. Ferrand, and H. Mariette, Quantum Structures of II–VI Diluted Magnetic Semiconductors Agnieszka Wolos and Maria Kaminska, Magnetic Impurities in Wide Band-gap III–V Semiconductors Tomasz Dietl, Exchange Interactions and Nanoscale Phase Separations in Magnetically Doped Semiconductors Hiroshi Katayama-Yoshida, Kazunori Sato, Tetsuya Fukushima, Masayuki Toyoda, Hidetoshi Kizaki, and An van Dinh, Computational Nano-Materials Design for the Wide Band-Gap and High-TC Semiconductor Spintronics Masaaki Tanaka, Masafumi Yokoyama, Pham Nam Hai, and Shinobu Ohya, Properties and Functionalities of MnAs/III–V Hybrid and Composite Structures

Volume 83 Semiconductors and Semimetals

T. Scholak, F. Mintert, T. Wellens, and A. Buchleitner, Transport and Entanglement P. Nalbach and M. Thorwart, Quantum Coherence and Entanglement in Photosynthetic Light-Harvesting Complexes Richard J. Cogdell and Jurgen€ Köhler, Sunlight, Purple Bacteria, and Quantum Mechanics: How Purple Bacteria Harness Quantum Mechanics for Efficient Light Harvesting

Volume 84 Semiconductors and Semimetals David Z.-Y. Ting, Alexander Soibel, Linda Höglund, Jean Nguyen, Cory J. Hill, Arezou Khoshakhlagh, and Sarath D. Gunapala, Type-II Superlattice Infrared Detectors S. D. Gunapala, S. V. Bandara, S. B. Rafol, and D. Z. Ting, QuantumWell Infrared Photodetectors Ajit V. Barve and Sanjay Krishna, Quantum Dot Infrared Photodetectors J. C. Cao and H. C. Liu, Terahertz Semiconductor Quantum Well Photodetectors A. G. U. Perera, Homo- and Heterojunction InterfacialWorkfunction Internal Photo-Emission Detectors from UV to IR David R. Rhiger, HgCdTe Long-Wave Infrared Detectors

Volume 85 Semiconductors and Semimetals

Darius Abramavicius, Vytautas Butkus, and Leonas Valkunas, Interplay of Exciton Coherence and Dissipation in Molecular Aggregates Oliver Kuhn€ and Stefan Lochbrunner, Quantum Dynamics and Spectroscopy of Excitons in Molecular Aggregates Carsten Olbrich and Ulrich Kleinekathöfer, From Atomistic Modeling to Electronic Properties of Light- Harvesting Systems Alex W. Chin, Susana F. Huelga, and Martin B. Plenio, Chain Representations of Open Quantum Systems and Their Numerical Simulation with Time-Adaptive Density Matrix Renormalisation Group Methods Avinash Kolli and Alexandra Olaya-Castro, Electronic Excitation Dynamics in a Framework of Shifted Oscillators 166 Contents of Volumes in this Series

E. Lifshitz, R. Vaxenburg, G. I. Maikov, D. Yanover, A. Brusilovski, J. Tilchin, and A. Sashchiuk, The Significance of Alloy Colloidal Quantum Dots Elizabeth von Hauff, The Role of Molecular Structure and Conformation in Polymer Electronics Koen Vandewal, Kristofer Tvingstedt, and Olle Ingana¨s, Charge Transfer States in Organic Donor–Acceptor Solar Cells Carsten Deibel, Photocurrent Generation in Organic Solar Cells

Volume 86 Advances in Semiconductor Lasers Joseph P. Donnelly, Paul W. Juodawlkis, Robin Huang, Jason J. Plant, Gary M. Smith, Leo J. Missaggia, William Loh, Shawn M. Redmond, Bien Chann, Michael K. Connors, Reuel B. Swint, and George W. Turner, High-Power Slab-Coupled Optical Waveguide Lasers and Amplifiers P. Crump, O. Brox, F. Bugge, J. Fricke, C. Schultz, M. Spreemann, B. Sumpf, H. Wenzel, and G. Erbert, High-Power, High-Efficiency Monolithic Edge-Emitting GaAs-Based Lasers with Narrow Spectral Widths E. A. Avrutin and E. U. Rafailov, Advances in Mode-Locked Semiconductor Lasers K. M. Kelchner, S. P. DenBaars, and J. S. Speck, GaN Laser Diodes on Nonpolar and Semipolar Planes Eric Tournie´ and Alexei N. Baranov, Mid-Infrared Semiconductor Lasers: A Review Dominic F. Siriani and Kent D. Choquette, Coherent Coupling of Vertical-Cavity Surface-Emitting Laser Arrays Anne C. Tropper, Adrian H. Quarterman, and Keith G. Wilcox, Ultrafast Vertical-External-Cavity Surface- Emitting Semiconductor Lasers Soon-Hong Kwon, Hong-Gyu Park, and Yong-Hee Lee, Photonic Crystal Lasers Martin T. Hill, Metallic and Plasmonic Nanolasers Mark T. Crowley, Nader A. Naderi, Hui Su, Frederic Grillot, and Luke F. Lester, GaAs-Based Quantum Dot Lasers Philip Poole, InP-Based Quantum Dot Lasers C. Z. Ning, Semiconductor Nanowire Lasers

Volume 87 Advances in Photovoltaics: Volume 1

Hans-Josef Fell, Foreword Eicke R. Weber and Gerhard P. Willeke, Introduction Gerhard P. Willeke and Armin Ra¨uber, On The History of Terrestrial PV Development: With a Focus on Germany Paula Mints, Overview of Photovoltaic Production, Markets, and Perspectives Gregory F. Nemet and Diana Husmann, PV Learning Curves and Cost Dynamics Martin A. Green, Photovoltaic Material Resources Laszlo Fabry and Karl Hesse, Crystalline Silicon Feedstock Preparation and Analysis

Volume 88 Oxide Semiconductors

John L. Lyons, Anderson Janotti, and Chris G. Van de Walle, Theory and Modeling of Oxide Semiconductors Contents of Volumes in this Series 167

Filip Tuomisto, Open Volume Defects: Positron Annihilation Spectroscopy Lasse Vines and Andrej Kuznetsov, Bulk Growth and Impurities Leonard J. Brillson, Surfaces and Interfaces of Zinc Oxide Tadatsugu Minami, Transparent Conductive Oxides for Transparent Electrode Applications Bruno K. Meyer, Angelika Polity, Daniel Reppin, Martin Becker, Philipp Hering, Benedikt Kramm, Peter J. Klar, Thomas Sander, Christian Reindl, Christian Heiliger, Markus Heinemann, Christian Muller,€ and Carsten

Ronning, The Physics of Copper Oxide (Cu2O) Cheng Song and Feng Pan, Transition Metal-Doped Magnetic Oxides Katharina Grossmann, Udo Weimar, and Nicolae Barsan, Semiconducting Metal Oxides Based Gas Sensors John F. Wager and Bao Yeh, Oxide Thin-Film Transistors: Device Physics

Volume 89 Advances in Photovoltaics: Part 2

Otwin Breitenstein, The Physics of Industrial Crystalline Silicon Solar Cells Matthias Heuer, Metallurgical Grade and Metallurgically Refined Silicon for Photovoltaics Harry Wirth, Crystalline Silicon PV Module Technology Ulf Blieske and Gunther Stollwerck, Glass and Other Encapsulation Materials Karsten Bothe and David Hinken, Quantitative Luminescence Characterization of Crystalline Silicon Solar Cells