Copper Diffusion Through Plated Nickel Barrier Layers for Laser-Doped Selective-Emitter Silicon Solar Cells

Copper Diffusion through Plated Nickel Barrier Layers for Laser-Doped Selective-Emitter Silicon Solar Cells

Shahla Zamani Javid

School of Photovoltaic and Renewable Energy Engineering

University of New South Wale Sydney, Australia

A thesis submitted to University of New South Wales in fulfilment of a requirement for the degree of Masters by Research

2012

ORIGINALITY STATEMENT

„I hereby declare that this submission is my own work and to the best of my knowledge it contains no materials previously published or written by another person, or substantial proportions of material which have been accepted for the award of any other degree or diploma at UNSW or any other educational institution, except where due acknowledgement is made in the thesis. Any contribution made to the research by others, with whom I have worked at UNSW or elsewhere, is explicitly acknowledged in the thesis. I also declare that the intellectual content of this thesis is the product of my own work, except to the extent that assistance from others in the project's design and conception or in style, presentation and linguistic expression is acknowledged.‟

Signed …………………Sh. Javid………………………….

Date ……………22nd of May 2012…………………......

iii

„I hereby grant the University of New South Wales or its agents the right to archive and to make available my thesis or dissertation in whole or part in the University libraries in all forms of media, now or here after known, subject to the provisions of the Copyright Act 1968. I retain all proprietary rights, such as patent rights. I also retain the right to use in future works (such as articles or books) all or part of this thesis or dissertation. I also authorise University Microfilms to use the 350 word abstract of my thesis in Dissertation Abstract International (this is applicable to doctoral theses only). I have either used no substantial portions of copyright material in my thesis or I have obtained permission to use copyright material; where permission has not been granted I have applied/will apply for a partial restriction of the digital copy of my thesis or dissertation.'

Signed …………………Sh. Javid………………………….

Date ……………22nd of May 2012…………………......

AUTHENTICITY STATEMENT

„I certify that the Library deposit digital copy is a direct equivalent of the final officially approved version of my thesis. No emendation of content has occurred and if there are any minor variations in formatting, they are the result of the conversion to digital format.‟

Signed …………………Sh. Javid………………………….

Date ……………22nd of May 2012…………………......

vi Abstract

Nickel barrier layers are commonly-used in metallisation of silicon-based microelectronic and photovoltaic devices to prevent copper diffusion from metal contacts to the underlying silicon. This thesis compares the effectiveness of nickel barrier layers, formed using: (i) electroless plating; and (ii) light-induced plating (LIP), in preventing copper diffusion into mono-crystalline silicon laser-doped selective emitter (LDSE) solar cells.

Energy dispersive X-ray spectroscopy was used in conjunction with transmission electron microscopy for chemical identification of copper-rich regions in silicon under the nickel/copper metal contacts of heat-treated cells. It was observed that copper can penetrate through both electroless-plated and LIP nickel barrier layers in LDSE cells that have been heated for 5 hours at 200 °C or 400 °C, followed by quenching in ethylene glycol (fast cooling).

Current-voltage measurements were used to characterise the electrical performance of LDSE cells with different nickel barrier layers. Significant degradations in cell efficiency, fill factor and open-circuit voltage were observed for LDSE cells with electroless-plated nickel barrier layers after heating at 200 °C and subsequent quenching in ethylene glycol. In the worst case, the fill factor was reduced from 66% to 31%, causing severe efficiency degradation. However, the fill factor of cells with LIP nickel barrier layers were only reduced from 66% to 58% under the same conditions. When cells were cooled slowly the electrical performance of the cells was only slightly affected and was independent of the type of nickel barrier layer. Furthermore, LDSE cells with electroless-plated nickel barrier layers that were originally heat-treated followed by fast cooling, were exposed to a second heat treatment at 250 °C for 2 hours, followed by slow cooling. This resulted in a significant improvement in the lifetime.

These results suggest that copper may diffuse through both electroless plated and LIP nickel barrier layers at temperatures as low as 200 °C. However, on slow cooling (which is likely in the field), the copper can out-diffuse to regions of the cell where the

vii cell‟s electrical performance is not degraded. These observations suggest that some commonly-accepted views on the ability of nickel barrier layers to prevent copper diffusion in silicon solar cells may need to be challenged and further investigated.

viii Acknowledgments

Firstly, I am most grateful to my supervisors Dr. Alison Lennon and Professor Stuart Wenham. Thanks Stuart for dragging me into the world of photovoltaic research and supporting me in critical stages throughout my degree. Also, thank you for providing me with opportunities to gain invaluable experience in the lab.

Special thanks to Alison for giving me the opportunity to work on this project. This has indeed been a humbling and life-changing experience. Thank you for your technical guidance, discussions about research direction, constant encouragement and reassurance, dedicated and timely proofreading despite your extremely busy schedule, being a great mentor and your cool “it‟s gonna be ok” attitude.

I would also like to express my appreciation to Dr. Richard Corkish for provision of my scholarship which assisted me with the completion of this degree. Thank you for all your help and support in the last few years.

Many thanks to Yael Augarten, for sharing with me the PL-based shunt measurement technique that she developed during her PhD. Thank you for all the lively PL conversations, enthusiastically answering my random PL questions (even during baking at home) and helping me make sense of the PL results presented in Chapter 4.

I would also like to thank all members of the growing first generation group at UNSW, in particular Martha Lenio, Roland Utama, Ziv Hameiri and Adeline Sugianto for initial training in the lab, processing tips and tricks and many discussions on solar cells at different stages of my research. Many thanks to Dr. Budi Tjahjono for technical discussions on solar cell fabrication and laser doping and for being a good friend. The rest of the team, Brett Hallam, Nicole Kuepper, Yu Yao, GuangQi Xu (Stanley) and Catherine Chan should also be thanked for their help along the way. Special thanks to my dear friends Danny Chen and Julie Kwan for their support and kindness and for being true friends in the last few years.

ix I am also grateful to Dr. Tom Puzzer for sharing with me his vast knowledge of electron microscopy and helping me master the microanalysis technique with TEM; Dr. Ivan Perez-Wurfl and Dr. Oliver Kunz for fruitful discussions on electrical measurement approach and analysis and Dr. Daniel Inns for reading and commenting on Chapter 2.

Thanks to the LDOT team, Dr. Nick Shaw, Kian Chin, Alan Yee, Bernhard Vogl (and Nancy), for making sure the labs run smoothly. Thanks also to Lawrence Soria for looking after our computers.

I am grateful to the EMU staff at UNSW, especially Sean Lim and Katie Levick, for enormous help with FIB, TEM sample preparation and TEM/EDX analysis presented in Chapter 3 and Eugene White for his assistance with SEM.

I would like to express my deepest gratitude to my lovely family; Mum and Dad for their endless love and support and my sisters Nasrin and Simin, for doing all the wonderful things that only sisters can do. Many thanks to Nasrin and Zia for always being there for me and my adorable nieces Melika, Ayla and Bahharr, for their unconditional love that kept me sane during this time.

And the most special thanks to my fiancé Chris for his love and support, careful proofreading and exotic cooking, late night technical discussions and helping me make sense of the weird electrical measurement results that eventually became Chapter 4. Thank you for sharing the best and the worst during this intense stage of my life and for being the most amazing “novio”. This thesis would have been impossible without you.

x List of Abbreviations

AES Auger Electron Spectroscopy ARC Anti-Reflection Coating BC Buried Contact BSF Back Surface Field CB Conduction Band CCD Charge-Coupled Device CFL Compact Fluorescent Light CVD Chemical Vapour Deposition CZ Czochralski DI Deionised DSBC Double-Sided Buried Contact EBIC Electron Beam Induced Current EDX Energy Dispersive X-ray Spectrometry EL Electroluminescence EWT Emitter Wrap-Through FF Fill Factor FIB Focused Ion Beam FWHM Full Width at Half Maximum IBBC Interdigitated Backside Buried Contact inSECT inline Selective-Emitter Cell Technology LCP Laser Chemical Processing LDSE Laser-Doped Selective-Emitter LIP Light-Induced Plating MPP Maximum Power Point MWT Metal Wrap-Through PECVD Plasma-Enhanced Chemical Vapour Deposition PERL Passivated Emitter Rear Locally-Diffused pFF pseudo Fill Factor

xi PL Photoluminescence PSG Phosphosilicate Glass PVD Physical Vapour Deposition PXRD Powder X-ray Diffraction rpm Revolutions per minute SCR Space Charge Region SEM Scanning Electron Microscope SIMS Secondary Ion Mass Spectrometry SRH Shockley-Read-Hall recombination Suntech Suntech Power, Co. Ltd TEM Transmission Electron Microscopy TID Transient Ion Drift TXRF Total Reflection X-ray Fluorescence UNSW University of New South Wales VB Valence Band XRF X-ray Fluorescence

xii List of Symbols

Ag Silver Al Aluminium at.% Atomic percent Au Gold B Boron C Carbon Co Cobalt Cu Copper

Cu3Si Copper silicide

CuSO4 Copper sulphate

EC Energy of conduction band eV Electron-volt Fe Iron Ga Gallium

H2O2 Hydrogen peroxide

H2SO4 Sulphuric acid

H3BO3 Boric acid

H3PO4 Phosphoric acid HF Hydrofluoric acid I Current

IL Light-generated current

Impp Current at maximum power point

Irsh Current flowing through the shunt path

Isc Short circuit current I-V Current-Voltage

Jo Saturation current density

JSC Short-circuit current density K/s Kelvin/second

KB Boltzmann constant m Local ideality factor

xiii m-V Local ideality factor-Voltage n Cell ideality factor

N2 Nitrogen gas

NH4OH Ammonium hydroxide Ni Nickel NiSi Nickel silicide O Oxygen P Phosphorus

Pin Incident solar power Pt Platinum q Electron charge R Reflectance

Rs Series Resistance

Rsh Shunt Resistance Si Silicon

SiNx Silicon nitride

SiO2 Silicon dioxide T Temperature Ta Tantalum Ti Titanium V Voltage

Vd Voltage across the p-n junction

Vmpp Voltage at maximum power point

Voc Open circuit voltage

VT Thermal Voltage w/o Weight/percent w/v Weight/Volume percent η Cell efficiency Ф Incident photon flux

xiv Table of Contents

Abstract...... v Acknowledgements...... vii List of Abbreviations...... ix List of Symbols...... xi Table of Contents...... xiii Chapter 1 Introduction ...... 1

1.1 Motivation ...... 1

1.2 Thesis Aim ...... 4

1.3 Scope of Work ...... 5

1.4 Thesis Outline ...... 5

Chapter 2 Literature Review: Nickel Barrier Layers for Copper Metallisation of Silicon ...... 7

2.1 Introduction ...... 7

2.2 Electroless Plating ...... 9

2.3 Light-Induced Plating ...... 10

2.4 Properties of Nickel Barrier Layer for Copper Diffusion ...... 12

2.5 Properties of Copper ...... 20

2.5.1 Defect Reactions of Copper in Silicon ...... 21

2.5.2 Out-Diffusion versus Precipitation ...... 24

2.6 Impact of Copper Contamination on p-n Junction Devices and Solar Cells .... 27

2.6.1 Influence of Copper on Minority Carrier Lifetime in Silicon ...... 27

2.6.2 Effect of Copper on Degradation of Solar Cell Efficiency ...... 28

2.7 Summary ...... 31

Chapter 3 Characterisation of Copper-rich Regions in Silicon Solar Cells ...... 33

3.1 Introduction ...... 33

3.1.1 Transmission Electron Microscopy (TEM) ...... 34

xv 3.1.2 Energy Dispersive X-ray Spectrometry (EDX) ...... 35

3.2 Verification of the Characterisation Method ...... 41

3.2.1 Experimental ...... 41

3.2.2 Results and Discussion ...... 46

3.2.3 Conclusion ...... 53

3.3 Comparison between Electroless Nickel and Light-Induced Plated Nickel as Diffusion Barrier Layers ...... 54

3.3.1 Experimental ...... 54

3.3.2 Results and Discussion ...... 59

3.4 EDX Analysis Artefacts and Common Problems ...... 70

3.5 Conclusion ...... 72

Chapter 4 Impact of Copper Contamination on the Electrical Properties of Laser- Doped Selective- Emitter Solar Cells ...... 75

4.1 Introduction ...... 75

4.2 Theory of Analysis Methods ...... 76

4.2.1 Characterisation of Current-Voltage Curves ...... 76

4.2.2 Photoluminescence Imaging ...... 82

4.3 Experimental ...... 94

4.4 Results ...... 95

4.4.1 Efficiency and I-V Characterisation ...... 95

4.4.2 Quantitative Current Measurements using Photoluminescence Imaging 107

4.5 Discussion ...... 110

4.5.1 The Influence of Plating Method ...... 110

4.5.2 The Influence of Cooling Conditions...... 114

4.6 Conclusion ...... 119

Chapter 5 Conclusions and Future Work...... 123

5.1 Contributions of this Thesis ...... 124

xvi 5.2 Future work ...... 125

5.3 Final Comments ...... 126

References ...... 129

xvii

Chapter 1 Introduction

1.1 Motivation

Solar photovoltaics is the conversion of sunlight directly into electricity by means of a solar cell device. Due to depletion of conventional energy resources, and prominent environmental issues such as global warming, the photovoltaic industry has become increasingly important as an alternative source of energy to fossil fuels. However, the cost of electricity generated by photovoltaics is still higher than that from conventional sources of power generation (Solarbuzz 2011). Consequently, one of the primary aims of solar cell research in recent decades is to reduce manufacturing costs, while maintaining or improving device efficiency.

In order to produce solar cells commercially, the fabrication processes need to be robust and low cost. Screen-printing, which was first used in the fabrication of silicon (Si) solar cells in the 1970‟s (Ralph 1975, Green 1995), is a mature and well established technology that still dominates the manufacturing of Si solar cells (Szlufcik et al. 1997, Hilali et al. 2004, Wenham et al. 2006b). However, this technology has some major drawbacks which limit cell efficiency. First, because of the width (over 100 µm) of the screen-printed fingers, significant shading losses can result. Second, the heavily- diffused phosphorus (P) emitter, that is required to reduce the contact resistance and enable wider spacing of the metal fingers, reduces the efficiency of collecting carriers generated by low-wavelength light. Finally, these cells typically have relatively high series resistance due to the lateral conduction of carriers in the emitter and resistive losses in the screen-printed metal grid (Wenham and Green 1996, Glunz 2007). Furthermore, the use of expensive silver (Ag) paste in the metallisation of the front contacts can increase the production cost of screen-printed solar cells (Green 2001, Bartsch et al. 2010, Firman et al. 2010).

1 In order to overcome these limitations, new solar cell structures have been investigated in the last 20 years in order to achieve higher cell efficiencies. Selective-emitter designs, which use heavily-doped regions under the metal contacts while leaving the remaining emitter lightly-doped, have been some of the more effective approaches to enhance solar cell efficiency (Hilali et al. 2004, Tjahjono et al. 2007b, Chunduri 2009). Several methods have been developed to pattern openings on the surface of solar cells in order to form selective-emitters. Some of the most notable are:

- Laser scribing, used in the fabrication of Buried Contact (BC) solar cells (Wenham 1988), Double-Sided Buried Contact (DSBC) solar cells (Ebong et al. 1996), Interdigitated Backside Buried Contact (IBBC) solar cells (Guo 2004);

- Laser drilling, used in the fabrication of Emitter Wrap-Through (EWT) (Gee et al. 1993, Schmit and Gee 2010) and Metal Wrap-Through (MWT) Si solar cells (Van Kerschaver et al. 1998, Fellmeth et al. 2010);

- Photolithography, used in the fabrication of high efficiency solar cells such as Passivated Emitter Rear Locally-diffused (PERL) (Zhao et al. 1996, Zhao et al. 1999) developed at the University of New South Wales (UNSW) and Stanford University‟s backside point-contact solar cell (Swanson et al. 1984, Sinton et al. 1989); and

- Direct and indirect inkjet patterning technologies, developed at UNSW, as lower-cost alternatives to photolithography (Utama et al. 2008, Lennon 2010).

All the abovementioned patterning methods facilitate a self-aligned metallisation scheme whereby metal is selectively deposited on the exposed heavily-doped regions, while the rest of the cell surface is covered by an anti-reflection coating (ARC) layer. On the other hand, these patterning methods are typically followed by two or more high temperature processes to form the heavily-doped selective-emitter. This further complicates the fabrication process and increases the manufacturing cost (Tjahjono et al. 2008, Chunduri 2009).

2 Approaches to form selective-emitters without the need for an additional high temperature process include:

- The inline selective-emitter cell technology (inSECT) method which uses inkjet printing to mask the contact regions and an additional etching step to etch back the remaining highly-doped area (Lauermann et al. 2009) and inFlexLine Plus developed at Centrotherm (Chunduri 2009). In both these technologies, aligned screen-printing is used to form metal contacts to the heavily-doped regions.

- Laser Chemical Processing (LCP) followed by light-induced plating (LIP) of Ni- Ag which also avoids the subsequent high temperature step and does not require masking to create the high doping regions underneath the front grid contact (Kray et al. 2008, Kray et al. 2010).

- Laser doping, developed at UNSW (Tjahjono et al. 2007a, Hameiri et al. 2008, Kray et al. 2008, Mai et al. 2009, Hameiri 2010b), is an effective way to create a selective-emitter that advantageously avoids the high temperature process. In this method, a laser beam is used to selectively remove the ARC to form openings, while simultaneously incorporating the dopants into the molten Si to create a heavily-doped region. In addition, this technology facilitates a self- aligned metallisation scheme to form metal contacts and is typically carried out by plating of nickel (Ni) and copper (Cu) (Tjahjono et al. 2010).

Metallisation of many of these selective-emitter technologies can be achieved by self- aligned metal plating of relatively inexpensive metals such as Ni and Cu. The main benefits of Cu are its low cost and high conductivity. However, it has high diffusivity and solubility in Si. Copper contamination in solar cells can lead to degradation of cell performance by introducing efficient minority carrier recombination sites (Istratov and Weber 1998, Bartsch et al. 2010, Hernandez et al. 2010). Therefore a barrier layer is required to prevent direct contact of Cu with the underlying Si in a solar cell device. A thin layer of Ni is typically used as a diffusion barrier and is deposited on the conductive surface on the contact areas, followed by Ni sintering to form a low resistance Ni silicide alloy at the contact interface (Olowolafe et al. 1976). Nickel silicide reduces the contact resistance between the metal and Si and improves the

3 adhesion of metal lines to Si (Wenham 1986). Copper is then plated on the sintered Ni to form the conductive bulk of the metal contacts.

Nickel-Cu contacts can be formed using electroless plating (Wenham 1986 , Mallory and Hajdu 1990) and more recently by LIP (Glunz et al. 2008, Yao 2009). However, the chemical structure of a Ni layer formed by electroless plating is different from that formed by electroplating and LIP (Rohan and O'Riordan 2003). This difference is largely due to the fact that in electroless plating other elements apart from Ni (e.g., P or boron (B)) can be co-deposited with the Ni and these elements alter the structural properties of the Ni layer. This means that the Ni layers resulting from the different deposition methods may differ in their abilities to act as diffusion barriers for Cu. Therefore, the effectiveness of different types of Ni barrier layers used in solar cell devices needs to be investigated.

1.2 Thesis Aim

The aim of this thesis was to analyse and compare the Cu diffusion barrier properties of Ni layers, deposited by electroless plating and LIP, in crystalline Si Laser-Doped Selective-Emitter (LDSE) solar cells. With the increasing use of Cu metallisation and LIP processes for solar cell devices, an improved understanding of the potential diffusion of Cu through Ni barrier layers is required if new metallisation processes, that are compatible with the long lifetime warranties of photovoltaic modules, are to be developed. It is well known that Cu can easily diffuse into Si when heated and degrade device performance (Istratov and Weber 2002). This does not only extend to heating processes during fabrication of the solar cells. Solar modules generate heat as they produce electricity and therefore there is a risk of Cu diffusing into the cell and subsequently degrading cell performance if barrier layers are insufficient. Furthermore, solar photovoltaics modules are frequently installed in areas with high ambient temperatures and therefore may be exposed to high temperatures for extended periods of time. Modules are typically guaranteed to have a lifetime of 25 years, with some manufacturers aiming for a 30-year lifespan (Wenham et al. 2006a). Thus, it is critical to make sure that metallisation techniques which use Cu for solar cell metallic contacts are compatible with the expected lifetime warranty of modules.

4 Another important motivation for this study is that some solar cell manufacturers are currently seeking new methods to separate wet chemical processing from dry processing steps so as to simplify the equipment design, minimise the fabrication process steps and consequently reduce the manufacturing costs. One way to achieve this is to perform the sintering step, after completing both Ni and Cu plating. However, this can increase the risk of Cu penetrating through the Ni layer into Si during the sintering process, thus reducing the process reliability. A more comprehensive understanding of the effects of diffusing Cu in high temperatures will assist the development of new metallisation processes. This study was undertaken to investigate potential problems of employing Ni/Cu metallisation schemes in solar cell device fabrication in an industrial environment.

1.3 Scope of Work

This thesis is focused on characterisation of single-crystalline Si LDSE solar cells which were intentionally contaminated with Cu by thermal diffusion. This study combines a direct method to detect penetration of Cu through electroless Ni and LIP Ni diffusion barriers into the Si substrate and a number of electrical measurement techniques to analyse the effect of Cu on the electrical performance of LDSE solar cells. Special attention is paid to the influence of Cu contamination on the formation of localised shunting and/or Schottky diode contacts and the increase in the severity of the shunt in Cu contaminated cells.

1.4 Thesis Outline

Chapter 2 begins with a review of research into the properties of different types of Ni barrier layers for Cu in the metallisation of Si electronic devices. Firstly, brief background information on electroless plating and LIP techniques is presented. The barrier properties of electroless Ni and light-induced plated Ni in preventing the diffusion of Cu into Si are then discussed. Previous studies suggest that the Ni-P compound resulting from electroless Ni plating is more effective as a Cu diffusion barrier layer than the pure Ni deposited by LIP. Subsequently, the progress made in the research of the physical behaviour of Cu in Si is reviewed with an emphasis on different

5 defect reaction paths of Cu in Si. Finally, the impact of Cu contamination on p-n junction devices and Si solar cells is reviewed.

Chapter 3 presents a study of chemical identification of Cu-rich clusters in Si. The chapter begins with brief background information of electron microscopy and basic principles of elemental analysis. To investigate the effectiveness of Ni barrier layers in preventing the penetration of Cu into Si, LDSE solar cells are fabricated with electroless Ni and LIP Ni as diffusion barrier layers and Cu plated to form the metallic contacts. The finished cells are exposed to thermal treatments of different temperatures and durations to intentionally diffuse Cu from the plated metallic contacts into the Si. The diffusion of Cu through different types of Ni barrier layers is analysed using energy dispersive X-ray spectroscopy (EDX), in conjunction with transmission electron microscope (TEM).

If Cu penetrates through the Ni barrier layers into the Si, it may form Cu precipitates which act as active recombination centres for minority carriers. Chapter 4 investigates the effects of Cu contamination on the electrical performance of LDSE solar cells. A number of device characterisation methods such as examination of the current-voltage (I-V) curves and photoluminescence (PL) imaging are employed to address the objectives of this chapter. The I-V analysis is used to qualitatively examine the sources of high-recombination regions and formation of Schottky diode contacts whereas the PL imaging is used to quantitatively determine the recombination current in the shunted regions of the cells. The discussion is focussed on two major aspects. The first part studies the influence of different plating methods on the electrical performance of solar cells. The second part investigates the effect of different cooling conditions after termination of heat treatment on the electrical properties of the cells.

Finally, in Chapter 5, the important results of this study and suggestions for possible future work are presented.

6 Chapter 2 Literature Review: Nickel Barrier Layers for Copper Metallisation of Silicon

2.1 Introduction

Metallisation is one of the most critical steps in the fabrication of solar cells because it enables the collection of the light generated current in the cell. Ideally the metallisation method is low cost and metal contacts have high conductivity, low contact resistance and adhere adequately to the Si surface (Fisher 2007, Lee 2009).

Metallisation of commercially produced Si solar cells has been dominated by screen- printing technology for the past thirty years (Ralph 1975) with screen-printed Ag pastes being used for front grid contacts and aluminium (Al) pastes being used for rear contacting. This metallisation technology is robust, well understood and supported by the availability of suitable production line equipment. However, it has limitations with regard to printing resolution as finger widths must typically be greater than 100 m (Mette et al. 2007). Furthermore, because screen-printing is a contact printing method, metallisation of thinner wafers can result in significant yield issues. Finally, Ag is an expensive metal and Ag pastes contribute a significant fraction to the cost of producing screen-printed solar cells (Green 2001, Bartsch et al. 2010, Firman et al. 2010).

Metallisation using metal plating has the potential to overcome these limitations, especially if cheaper metals such as Cu are used. In addition to its low cost, Cu has a much lower resistivity than screen-printed Ag with resistivity values reported as low as 1.7 μΩ cm (Cabral et al. 1992, Lee et al. 1998). The resistivity of Ag paste is reported to be around two times the resistivity of pure Ag (1.6 μΩ cm) for busbars and six times higher than that for fine features such as fingers, due to the porous structure of the paste (de Moor et al. 1997). Copper is however, a fast diffuser in Si and as a result, a barrier 7 layer is required to protect the underlying Si from Cu diffusion. Nickel has been widely used in microelectronic applications as well as solar cell devices to retard the diffusion of Cu into Si (Pinnel and Bennett 1976, Brady and Hovland 1981, Fisher 2007, Lee 2009).

Also, Ni silicide which is formed at the interface between Ni and Si by sintering the plated Ni has a low resistivity of 14 μΩ cm and reduces the contact resistance (Kim 2005). Sintering the cells after Ni plating also improves the adhesion of the metal contacts to the Si substrate. Copper can then be plated onto the Ni/ Ni silicide regions to form high quality contacts (Wenham 1986, O'Sullivan et al. 1998, Fisher 2007).

Plated Ni/Cu contacts have been used extensively in the fabrication of Si solar cells in the laboratory (Wenham 1986, Wenham and Green 1988a, Wenham and Green 1988b, Kim 2005, Tjahjono et al. 2007b, Riegel et al. 2008). In some cases the metal was plated using electroplating processes which involve forming electrical contacts to the cell to enable the plating (Lee et al. 2002), though most uses of Ni/Cu plating have involved the use of electroless plating and more recently LIP (Glunz et al. 2008, Tjahjono et al. 2010). Although electroless plating is attractive in that it does not require electrical contact to be made to the cell, it suffers from a low deposition rate and a strict requirement to closely control the pH, temperature and chemical composition of the plating bath. Light-induced plating is similar to electroplating, except that the solar cell‟s current generating capability is utilised in the plating process instead of an external power supply, thus eliminating the need for external electrodes and contacts to the cell (Durkee 1979). This reduces cost and simplifies the plating process and so is considered as a more promising alternative for industrial environments. Furthermore, LIP is a low temperature process and results in faster metal deposition than electroless plating (Mette et al. 2006, Yao 2009).

However, to date metallisation by plating has only been used in several commercially produced solar cells. Examples include the BP Solar Saturn Si solar cells (Bruton 1992, Mason 2004) and the Suntech Power, Co. Ltd Pluto cells (EnergyMatters 2009, Shi et al. 2009).

8 With the increased use of Cu contacts for Si solar cells, whether by electroless plating or LIP, solar cell devices are more exposed to the risk of being contaminated with Cu. Therefore, it is important to understand the effectiveness of different types of Ni barrier layers in preventing Cu diffusion into the underlying Si. Nickel barrier layers deposited by electroless plating and LIP differ in chemical structure and morphology (Turn and Owen 1974, O'Sullivan et al. 1998, Rohan and O'Riordan 2003) and so may vary in their ability to block the diffusion of Cu into the cell. Furthermore, if barrier layer quality is to be investigated by detection of Cu in Si then it is critical to also understand the physical and electrical properties of Cu in Si.

An overview of electroless plating and LIP techniques is provided in Section 2.2‎ and Section 2.3‎ , respectively. Then, in Section ‎2.4, a review of previous research on the barrier layer properties of different types of Ni layers is provided. Finally, in Section 2.5‎ and Section ‎2.6 the physical and electrical properties of Cu in Si are reviewed and the effects of Cu contamination in solar cell devices are discussed.

2.2 Electroless Plating

Electroless plating is a simple, low cost and self-aligning metallisation technique. It is based on the selective deposition of metal onto a conductive substrate surface by immersing the substrate into a plating solution containing reducing agents and metal ions. Electroless plating does not require an external bias potential and is the result of localised electrochemical reactions.

Electroless plating of Ni (e.g., for barrier layers) can be achieved using a number of different commercially-available plating solutions. Plating formulations can vary in the type of reducing agent that they use. A commonly-used electroless Ni plating solution, and also the solution used in this work is ammonium-based with sodium hypophosphite

(NaH2PO2.H2O) as the reducing agent. Ammonium hydroxide (NH4OH) is used to control the pH of the bath (Guo and Cotter 2005). According to the following equations, the electrons are generated by the oxidation reaction of sodium hypophosphite and the Ni ions present in the plating solution are converted to solid Ni which is deposited on the conductive surface (Guo and Cotter 2005):

9 − − + − 퐻2푃푂2 + 퐻2푂 → 퐻2푃푂3 + 2퐻 + 2푒 (2.1)

2+ − 푁𝑖 + 2푒 → 푁𝑖 (2.2)

An anti-reflection coating comprising Si dioxide (SiO2) or Si nitride (SiNx) acts as a plating mask and the Ni ions plate only on the conductive surface of the contact area. In the experiments reported in this thesis the contact areas are heavily-doped regions formed by laser doping.

Electroless plating u tilising sodium hypophosphite as the reducing agent results in co- deposition of P and Ni. The P content in the deposited Ni-P alloy depends on the pH and temperature of the bath, and the hypophosphite concentration (O'Sullivan et al. 1998, Fisher 2007). As will be explained in Section ‎2.4, the P present in the Ni film deposited by electroless plating makes a considerable difference in the quality of the diffusion barrier as compared to the pure Ni layer deposited using electroplating or LIP.

2.3 Light-Induced Plating

Light-induced plating utilizes the same principles of electroplating. Table 2.1 outlines the five essential components in electroplating and the equivalent components in LIP methods.

In electroplating, both the anode and cathode are immersed in the electrolyte solution. The direct current flow to the anode results in the oxidation of the anode metal (i.e., loss of electrons) and the formation of positively charged metal ions (cations) in the electrolyte. These cations are attracted to the negatively charged cathode where they accept electrons and are reduced to form metal on the cathode.

10 Table 2.1: A comparison between the elemental components of electroplating and LIP systems.

Electroplating LIP

A power supply to drive the deposition A solar cell acting as a power supply. process.

An anode comprising either the rear p-type An anode usually made of the metal to be electrode of the solar cell or a metal block in plated. contact with the p-type Si of the solar cell.

A cathode typically a conductive metal and is One or more regions of n-type Si of the solar the piece to be plated. cell acting as the cathode.

Electrolyte (the bath solution containing Electrolyte (the bath solution containing metal ions). metal ions).

Electric wires used to connect the positive Silicon of solar cell conducting current terminal of the power supply to the anode and between the anode and cathode. the negative terminal to the cathode.

In LIP technique, a light source is used to generate electron-hole pairs in the solar cell which is immersed in the plating bath. In some implementations of LIP, the Al rear electrode of the solar cell which is in contact with the p-type Si, acts as an anode (see Figure 2‎ .1). The holes collected in this region will cause the rear surface Al to be oxidised. The Al loses electrons and the metal ions (Al3+) enter the plating solution (oxidation process). Meanwhile, light-generated electrons are collected by the n-type Si of the cell where they are accepted by the positively charged metal ions on the exposed n-type contact regions on the illuminated surface of the solar cell (reduction process). The electric circuit is therefore completed and the metal is deposited on the required area.

11 Metal ions Cathodic (e.g., Cu2+) reaction (reduction)

Anodic reaction (oxidation) Al3+

Figure 2.1: A schematic diagram to illustrate the process of anodic oxidation and cathodic reduction reactions during the LIP process (Used with the permission of Tjahjono 2010).

2.4 Properties of Nickel Barrier Layer for Copper Diffusion

Copper metal has several advantages such as low cost, low electrical resistivity and high electromigration resistance. Consequently, Cu is generally used in the metallisation of interconnects in microelectronic applications and top surface metal contacts in solar cell devices (Baumann et al. 1997, Lee et al. 1998, Rohan et al. 2002, Liu et al. 2010). Electroless Cu plating is the preferred deposition method when compared to Chemical Vapour Deposition (CVD) and sputtering because it is a simple and low cost process and metal is selectivity deposited only on the catalytic surface of the substrate (Lee et al. 1998). However, Cu is a fast diffuser in Si even at low temperature, and its solubility in Si is very high at high temperatures. Copper, if diffused into the solar cells will severely degrade the device performance (Buonassisi et al. 2005, Bartsch et al. 2010). Therefore, a highly reliable metallic diffusion barrier between the metal layer and Si substrate is required to protect the Si from Cu diffusion. The essential characteristics of a good diffusion barrier layer include uniform coverage, low electrical resistivity, strong adhesion and slow transport rates to the neighbouring materials (Lee et al. 1998, Liu et al. 2010).

Turn and Owen investigated the ability of several plated metals and alloys to prevent the interdiffusion of Cu and Au in 1974 (Turn and Owen 1974). They examined 12 electrodeposited and electroless deposited Ni barriers with different thicknesses onto a gold (Au) substrate and compared the interdiffusion results of these samples with the results on a sample which had no barrier. Using an electron microprobe, the concentration-distance profiles of Cu, Au and Ni were measured. A comparison of Cu penetration values obtained at high temperatures (400 °C and 550 °C) for 12 to 53 hours is shown in Table ‎2.2.

According to Table ‎2.2, both Ni alloys of 8 percent and 2-3 percent P effectively reduce the Cu penetration as compared to that of the sample in which there was no barrier. Also evident from this table, is the effect of P content of the barrier alloy in slowing Cu diffusion. The values of Cu penetration for electroless Ni alloy of 8 percent P are significantly lower than the Cu penetration for pure Ni deposited by electroplating. Similarly, the 2 μm thick of the low P deposit of Ni (Ni- 2-3 percent P) has blocked or considerably slowed down the Cu diffusion compared to the pure Ni layer. Thus, it was concluded that an electrolessly deposited Ni-P alloy is the most effective metallic barrier in preventing the interdiffusion of Cu and Au (Turn and Owen 1974, Lee et al. 1998, O'Sullivan et al. 1998). Furthermore it was found that an increase in the P content of the Ni-P alloy improves the effectiveness of the electroless Ni deposit to act as a barrier and thus decrease the diffusion of Cu (Turn and Owen 1974, O'Sullivan et al. 1998, Rohan and O'Riordan 2003).

13 Table 2.2: Penetration of Cu through different diffusion barriers into Au (Turn and Owen 1974).

Heat Treatment Parameters Barrier Layer Barrier Layer 400°C, 24 hr. 400°C, 53 hr. 550°C, 12 hr. Thickness (µm) Copper penetration values (µm)

No barrier - 7 13 18

Electroless Ni 0.25 1 12 18 (Ni-8 percent P) Electroless Ni 0.5 1 6 15 (Ni-8 percent P) Electroless Ni 1 1 1 8 (Ni-8 percent P) Electroless Ni 2 No diffusion No diffusion No diffusion (Ni-8 percent P) Electro-plated Ni 1 1 6 10 (pure Ni) Electro-plated Ni 2 5 12 12 (pure Ni) Electroless Ni 1 No diffusion 2 10 (Ni-2-3 percent P) Electroless Ni 2 No diffusion No diffusion 6 (Ni-2-3 percent P)

However, this study did not present Cu diffusion data for lower temperatures. Furthermore, it is not intuitive that Cu diffusion could increase with increasing thickness of the electroplated Ni barrier layer. This apparent anomaly was not explained in the study.

Pinnel and Bennett in 1976 proposed a mechanism for the interdiffusion of Cu and Au through a Ni barrier layer (Pinnel and Bennett 1976). They prepared Cu/Ni/Au samples using electroplating to deposit the Ni. An electron microprobe was used to measure the penetration curves over a temperature range of 150 to 750 °C. It was demonstrated that annealing the Cu/Ni/Au samples at a temperature of 500 °C or below, causes Au to diffuse through the Ni layer and form a concentration build-up at the Cu/Ni interface. Similarly, Cu penetrates through the Ni and accumulates at the Au/Ni interface. It was also shown that Cu and Au diffuse through grain boundaries and dislocations in the Ni layer without entering the Ni lattice (Turn and Owen 1974, Pinnel and Bennett 1976). In

14 fact, the Ni layer acts more as a transit passage layer. Copper and Au atoms are prevented from entering the Ni lattice by solubility restrictions, but they appear to easily diffuse to the opposing surfaces. This is apparent from Figure 2.2, where there are Au and Cu peaks at the opposing interfaces respectively (Cu/Ni and Au/Ni), while no or very small amounts of Au and Cu are detected in the Ni layer (Pinnel and Bennett 1976). According to Pinnel and Bennett, although electroplated Ni restricts Cu penetration, it does not totally block the diffusion of Cu atoms to the Au layer.

Figure 2.2: Representative interdiffusion behaviour for Au/Ni/Cu layer stacks at an annealing temperature of 400 °C for 80 days (Pinnel and Bennett 1976). In this figure X-ray intensities of Au-Mα, Ni-Kα and Cu-Kα are measured with an electron microprobe, then overlaid on a SEM image of the sample.

Contrary to the results of Turn and Owen, it was also showed that Cu penetration decreases as the thickness of the Ni layer is increased (Pinnel and Bennett 1976, O'Sullivan et al. 1998).

Other authors have also supported the effectiveness of electroless Ni-P alloys as diffusion barriers over the pure Ni deposited by electroplating (Graham et al. 1965, O'Sullivan et al. 1998). According to a study by O‟Sullivan et. al., electrodeposited Ni, either from Ni sulphate or Ni sulfamate solutions are not effective in preventing Cu

15 diffusion. Auger electron spectroscopy (AES) analysis combined with sputter-profiling techniques was used to determine Cu interdiffusion in the Cu/barrier region. The Cu concentration (at.%), was measured as a function of sputter-removal time after annealing at 400 °C for 14 hours (O'Sullivan et al. 1998). Typically, the effectiveness of electrolessly deposited alloys in retarding Cu diffusion is attributed to the difference in the structure of the material and particularly the formation of intermetallic phosphides in the alloy. The Ni-P compound is reported to have an amorphous structure and the P is precipitated along the grain boundaries; thus blocking the diffusivity path through the material. The Ni-P compound is formed by heating to 350 °C which contributes to the effectiveness of the barrier properties (O'Sullivan et al. 1998, Rohan and O'Riordan 2003).

Another thorough study carried out by Lee et. al. in 1998 further substantiates the role of electroless Ni as a diffusion barrier (Lee et al. 1998). A multilayer structure of

Cu/Ni/poly-Si/SiO2 thin films was made using standard photolithography and electroless deposition of Ni and Cu, and heat treated at 300 °C for 30 minutes. The cross-sectional transmission electron microscopy (XTEM) and energy dispersive spectroscopy (EDS) analysis of the heat treated wafer is shown in Figure 2.3. The presence of Ni silicide is visible from the peak positions of Ni and Si in Figure 2.3c. As evidenced by the EDS analysis shown in Figure 2.3b-d, no Cu diffusion into the silicide layer was detected. The TEM micrograph of the cross-section of the electroless Ni/Si sample after being annealed at 350 °C for 1 hour is shown in Figure 2.4. As can be seen, the mono silicide of Ni (NiSi) has a columnar structure with no voids which blocks the Cu atoms from diffusing into this layer ( Lee et al. 1998).

According to this study, electroless Ni can effectively block the diffusion of Cu after annealing at 350 °C for one hour although poly-Si could diffuse through the NiSi layer. It was found that the electroless Ni layer has an amorphous structure until after sintering at 300 °C for 30 minutes in this study (Lee et al. 1998). The effectiveness of electrolessly plated Ni as a diffusion barrier is again attributed to its amorphous structure. Atomic transport is mainly dominant along grain boundaries and dislocations in crystalline materials (Pinnel and Bennett 1976, Lee et al. 1998, Liu et al. 2010) and P precipitation along the grain boundaries obstruct the diffusion paths.

16 This study however, does not extend the experiments to longer sintering times nor higher sintering temperatures. Also lacking is analysis of the performance of pure Ni, deposited by electroplating or LIP, as a diffusion barrier.

Figure 2.3: a) Cross-sectional TEM micrograph of a Cu/Ni/poly-Si/SiO2 film after a one hour heat treatment at 350 °C, b) EDS analysis of point A, c) EDS analysis of point B, d) EDS analysis of point C (Lee et al. 1998).

Figure 2.4: Cross-sectional TEM micrograph of a NiSi layer (Lee et al. 1998)

In 2003, Rohan et. al. characterised the crystallisation structure of electroless Ni with different P concentrations and different annealing temperatures (Rohan and O'Riordan 2003). Using powder X-ray diffraction (PXRD), they determined the amorphous structure of as-deposited electroless Ni-P alloys. They found that the amorphous structure of Ni-P changes to crystalline Ni and Ni3P with increasing temperature. This is apparent from Figure 2.5 which shows the X-ray diffraction pattern for a sample with 7.42 w/o P, as-plated and heated up to 400 °C.

Figure 2.5: PXRD of Ni-P (7.42% P) deposits as-plated and heated to 300, 350 and 400 °C (Rohan and O'Riordan 2003).

On the contrary, the degree of crystallinity in Ni-P deposits decreases as the P content is increased. Also, the grain sizes in general increase with increasing heat treatment temperature and decrease with increasing P content in the Ni-P alloy (Rohan and O'Riordan 2003).

It seems that formation of the Ni3P phase is completed after annealing at temperatures up to 400 °C in nitrogen (N2) environment. The EDX analysis of samples after heat treatment at 200 °C, showed no interdiffusion between Ni and Cu. However, annealing at 400 °C caused Cu diffusion into the Ni layer. It is therefore concluded that the amorphous structure of the Ni-P deposit improves its effectiveness as a barrier layer. Copper diffusion may occur if samples are heat treated at temperatures higher than 400 °C because annealing at this temperature transforms the structure of the deposit mostly to crystalline form. Fortunately, P tends to be uniformly distributed along the grain boundaries and therefore blocks the diffusion path through this layer (O'Sullivan et al. 1998, Rohan and O'Riordan 2003).

19 2.5 Properties of Copper

It is well established that transition metals such as Cu can dramatically degrade the electrical performance of Si devices (Weber 1983, Itsumi et al. 1997, Istratov and Weber 1998, Seibt et al. 1998, Buonassisi et al. 2005, Bartsch et al. 2010, Hernandez et al. 2010). The behaviour of Cu in Si is however different from other transition metals like iron (Fe) and Ni. Many studies have been performed to provide a better understanding of the electrical properties of Cu and its behaviour in Si (Shabani et al. 1996, Baumann et al. 1997, Istratov et al. 2000a, Istratov et al. 2000b, Istratov and Weber 2002).

According to their studies, Cu is a shallow donor in Si and so it diffuses in Si as a positively charged ion (Cu+). For this reason, the diffusion coefficient of Cu is not only dependant on temperature but it is also affected by the dopant type and concentration of the dopant atoms (conductivity and doping level) in Si. This means that the diffusivity of Cu in p-type Si is reduced because of the attraction between the positively charged Cu+ ions and negatively charged boron (B-). This pairing effect causes some of the Cu+ ions to be temporarily trapped by the immobile acceptors, while only a fraction of Cu+ ions are mobile in the Si wafer. Another point to consider is that this trapping is only temporary and the trapped Cu+ ions can be released again and become free to diffuse in Si. Also, as will be explained in Section 2.5.1, only a small number of trapped Cu+ ions make chemical bonds with B to form CuB pairs (Istratov et al. 2000a).

The diffusion coefficient of Cu in intrinsic Si is calculated as (Istratov et al. 2000a, Istratov and Weber 2002):

−4 0.18 ± 0.01 eV 2 퐷𝑖푛푡 = 3.0 ± 0.3 × 10 × 푒푥푝 − 푐푚 푠 (2.3) kB푇

where KB is the Boltzmann constant, the temperature T is measured in Kelvin. The effective diffusion coefficient is given by the following Equation for moderately-doped 17 -3 B (Na ≤ 10 cm ). In this Equation, the diffusion coefficient is reduced by the effect of trapping of mobile Cu+ ions by the immobile acceptor atoms (Istratov et al. 2000a, Istratov and Weber 2002):

20 −4 3 × 10 × 푒푥푝 −2090 푇 2 퐷푒푓푓 = −20 푐푚 푠 (2.4) 1 + 2.584 × 10 × 푒푥푝 4990 푇 × 푁푎 푇

Based on Equation 2.3, the intrinsic diffusion coefficient of Cu in Si at room temperature reaches 2.8 × 10-7 cm2/s which is high enough for Cu atoms to diffuse through a 600 µm intrinsic Si wafer in about an hour. According to Equation 2.4, the diffusivity decreases with increasing B acceptor concentration and increases with increasing temperature.

2.5.1 Defect Reactions of Copper in Silicon

As discussed in Section 2.5, Cu is very mobile in Si even at room temperature and therefore it will diffuse towards stable sinks within a Si wafer. The behaviour of Cu and its reaction mechanisms in Si are further studied by Istratovet.et. al. (Istratov et al. 2000a, Istratov and Weber 2002). According to (Graff 1995), the solubility of Cu in Si is given by:

23 −3 푆 = 5.51 × 10 × 푒푥푝 −1.49 푒푉 푘퐵푇 푐푚 (2.5) which gives a solubility of 1018 cm-3 at 1000 °C, and a very low value of less than 1 cm-3 Cu atoms at room temperature. Therefore, most of the interstitial Cu dissolved in the bulk of a Si wafer during annealing will take one of the following reaction paths after the Si wafer is cooled and returned to room temperature:

 Formation of point defects such as CuB pairs in the bulk. Copper and B can react to form CuB complexes, however as explained in Section 2.5, the high diffusivity of Cu in Si makes CuB pairs unstable even at room temperature. This is because the dissociation energy of CuB pairs is only 0.7 eV which is very low compared to the dissociation energy of 1.2 eV for FeB pairs for example. This causes CuB pairs, as well as other Cu point defects and complexes, to dissociate easily in Si at room temperature. A similar dissociation for FeB pairs takes place at 200 °C. Copper, as a result, diffuses towards more stable sinks such as extended defects in the bulk or preferably the surface of the wafer. Accordingly,

21 the concentration of electrically active point defects of Cu in the bulk of the Si wafer will be small, about 0.1% of the total Cu concentration in the bulk (Flink et al. 2000, Istratov et al. 2000a, Istratov and Weber 2002).

 Precipitation as Cu silicide (Cu3Si) in the bulk. It was observed in a study by Istratov et.al. in 1998 that Cu precipitates in Si may have different charge states and are dependent on the Fermi level in the sample. The reason is that the Cu precipitates form band like states in the upper half of the Si bandgap and have

the critical value of EC– 0.2 eV. In p-type Si, the Fermi level is below this limit which makes the Cu precipitates positively charged. As a result, positively- charged Cu ions are repelled and so the Cu precipitation is suppressed. On the other hand, when the Fermi level is raised above this value (in n-type Si for example), the Cu precipitates are neutral or negatively charged and accordingly, an electrostatic attraction will be created between Cu precipitates and the positively charged Cu ions. This will increase the precipitation of Cu in Si. Therefore, precipitation of Cu is more probable to take place in n-type Si than in p-type (Istratov et al. 1998, Flink et al. 2000, Istratov et al. 2000a). This is illustrated in Figure 2.6.

Another point to consider is that the concentration of interstitial Cu+ in Si affects the position of the Fermi level, in a similar way that dopant atoms like B or P do (Istratov et al. 2000b). Thus, during the precipitation of Cu, the concentration of interstitial Cu in the bulk decreases which leads to the Fermi level to decrease below the critical value. Consequently, the Cu precipitates become positively charged again and the precipitation rate will be reduced and out-diffusion will dominate (Flink et al. 2000, Istratov et al. 2000a, Istratov et al. 2000b, Istratov and Weber 2002). The out-diffusion behaviour of Cu is further discussed in Section 2.5.2.

Also, it was found that slow cooling of samples with a high initial Cu concentration leads to formation of large Cu silicide precipitates with low densities, mostly close to the wafer surface. These precipitates were grown mainly on pre-existing dislocation loops and stacking faults punched out by the existing precipitates. The size of these precipitates depends on the Cu 22 contamination level and cooling rate (Seibt et al. 1999). However, if samples were rapidly quenched, a high density up to 1013 cm-3 of small (~ 30-200 nm) Cu precipitates were distributed uniformly through the bulk of the wafer (Istratov and Weber 2002).

Figure 2.6: The effect of the electrostatic interaction between positively charged Cu ions and Cu precipitates on Cu precipitation in p and n-type Si (Istratov et al., 2000a, Istratov et al., 2000b)

 Formation of precipitates at extended defects in the Si wafer such as stacking faults, dislocations and grain boundaries (Istratov and Weber 2002). The fact that these extended defects provide stable sinks for metal impurities is however beneficial in wafer fabrication. It is the basis of the technique known as gettering, to detect and remove metal contamination.

 Out-diffusion to the surface of the wafer. Apparently, out-diffused Cu does not form chemical bonds with Si because it can be removed using a dilute

solution of hydrofluoric acid and hydrogen peroxide (HF: H2O2), from the surface of the p-type Si wafer. In this case, it was suggested that Cu

agglomerates at the Si/SiO2 interface or at the surface of the native oxide, or is incorporated into the native oxide (Shabani et al. 1996, Istratov et al. 2000a).

Also, observation of increasing growth of SiO2 in the presence of Cu

23 contamination indicates that Cu contributes to oxidation of Si (Istratov et al. 2000a).

As explained previously, the precipitation process of Cu is explained by an electrostatic model where the Fermi level position in the sample determines the precipitation or out-diffusion behaviour of Cu (Flink et al. 2000). Since the effective diffusion coefficient of Cu in Si reduces as the B doping level increases, out-diffusion of Cu will be slower in heavily B-doped wafers than in lightly-doped samples. Moreover, an increase in temperature will increase the rate of out-diffusion. However, it is reported that Cu is more stable in n-type Si. The reason may be the electrostatic attraction between positively charged Cu ions and the negatively charged Cu precipitates in n-type Si. Another possible reason reported in the literature is that chemical compounds may be formed between Cu and P (Shabani et al. 1996). It is observed that Cu mostly remains in the bulk of the wafer at room temperature but out-diffuses to the surface if it is sintered to 400 °C (Shabani et al. 1996, Istratov and Weber 2002).

 Segregation in p+area. This is used as a gettering technique to segregate Cu in the p+-Si substrate (Istratov et al. 2000a, Istratov and Weber 2002).

2.5.2 Out-Diffusion versus Precipitation

The process of Cu out-diffusion from n and p-type Si to wafer surfaces was studied by Shabani et.al. (Shabani et al. 1996). It was found that out-diffusion of Cu from B-doped p-type Si wafer could start at room temperature. The reason is attributed to instability of CuB pairs at room temperature which makes it easy for the Cu to migrate into the Si lattice and diffuse out to the wafer surface. It should be noted that the out-diffusion of Cu in p-type Si at such low temperatures will occur only if the surface oxide layer is removed. Conversely, out-diffusion of Cu in n-type Si is reported to take place at high temperature of about 400 °C. This is possibly due to the fact that pairs of P and Cu

(CunPm) in n-type Si are stable at room temperature and are broken at temperatures higher than 250 °C (Shabani et al. 1996, Istratov and Weber 2002).

24 The main two reaction paths of Cu in Si, i.e., out-diffusion and precipitation, have been explained thoroughly by Flink et.al. in 2000 (Flink et al. 2000). They demonstrated that in samples with high Cu concentration of about 1017-1018 cm-3, most of the Cu will precipitate in the bulk upon cooling. In contrast, for p-type Si with a low level of Cu contamination around 1014-1015 cm-3, total reflection X-ray fluorescence (TXRF) studies showed a complete out-diffusion of Cu to the surface, for both slow cooling and quenching. Also, the lower the doping level, the faster will be the out-diffusion (Flink et al. 2000).

Figure 2.7 shows the transient ion drift (TID) measurements of interstitial Cu concentration versus the Cu solubility at different in-diffusion temperatures. According to this figure, in samples with low Cu contamination, the concentration of interstitial Cu increases with Cu solubility until it reaches a critical contamination level. Above this level, the interstitial Cu concentration is decreased with increasing Cu solubility. This critical level is estimated to be equal to the B doping level plus 1016 cm-3. The synchrotron based X-ray fluorescence (XRF) method was used to determine the preferred reaction path of Cu above the critical level and the results are shown in Figure 2.8. As can be seen, if initial Cu contamination is larger than the critical level, the amount of precipitated Cu matches with Cu solubility. This means that most of the Cu will precipitate in the bulk. Whereas for the wafers with initial Cu contamination less than the critical level, out-diffusion is the main reaction path. This is shown in Figure 2.8 for the sample with high B doping of 2 × 1016 cm-3. There is no precipitation detected for this sample where the initial Cu contamination is less than the critical limit of 3 × 1016 cm-3 (Flink et al. 2000).

Figure 2.7: Interstitial Cu concentration versus Cu solubility at in-diffusion temperatures, measured by TID 30 minutes after quench at room temperature (Flink et al., 2000).

Figure 2.8: The concentration of precipitated Cu versus Cu solubility at in-diffusion temperatures (Flink et al., 2000).

26 2.6 Impact of Copper Contamination on p-n Junction Devices and Solar Cells

2.6.1 Influence of Copper on Minority Carrier Lifetime in Silicon

It was shown by several research groups that the minority carrier lifetime in Si wafers is dependent on the Cu contamination level (Naito and Nakashizu 1992, Istratov and Weber 1998, Istratov and Weber 2002 ). Figure 2.9 demonstrates the dependence of the lifetime in Czochralski (CZ) Si wafers on the surface concentration of Cu, Ni and Fe (Naito and Nakashizu 1992). It is clear from this figure that the level of Cu contamination has a stronger impact on lifetime in n-type Si than in p-type Si (Istratov and Weber 1998, Istratov and Weber 2002).

In n-type Si, the lifetime decreases as the Cu concentration level increases; whereas in p-type Si, the lifetime is shown to improve up to the Cu contamination level of 1013 cm-3 (Naito and Nakashizu 1992). However, in samples with a Cu surface concentration above 1014 cm-3, the lifetime decreases dramatically with increasing Cu contamination (Istratov and Weber 1998). The lifetime improvement observed in p-type Si with a low Cu contamination suggests that Cu has a passivating effect similar to that of hydrogen (Lee and Morrison 1988, Istratov and Weber 1998). This effect can be explained by formation of Cu complexes at active defects or impurity recombination sites, resulting in complexes with reduced recombination activity (Istratov and Weber 1998, Istratov and Weber 2002).

Figure 2.9: Dependence of normalised minority carrier lifetime in CZ Si on surface concentration of Cu, Ni and Fe. Identical thermal treatment followed by a slow cool was used for all surface contaminated samples (Naito and Nakashizu 1992).

2.6.2 Effect of Copper on Degradation of Solar Cell Efficiency

The effect of transition metals on the efficiency of solar cells was studied by Davis, Hopkins and Rohatgi (Davis et al. 1980, Hopkins and Rohatgi 1986). Figure 2.10 represents the results of Davis, Hopkins and Rohatgi for the threshold impurity concentrations for solar cell performance reduction (Istratov and Weber 1998). According to their studies, Cu contamination up to 1016 cm-3 does not affect the efficiency of CZ solar cells. Highly mobile Cu tends to form precipitates in Si where there are favourable nucleation sites or defects present in the material. Consequently, the Cu threshold concentration for adverse effects on multi-crystalline solar cell efficiency would be much lower than the threshold shown in this graph as they contain numerous grain boundaries, microdefects and dislocations, favouring the precipitation of Cu (Istratov and Weber 1998).

28 It was shown that interstitially dissolved Cu in Si does not greatly reduce the minority carrier lifetime since interstitial Cu can be gettered in non-critical regions of the device (Istratov and Weber 1998, Buonassisi et al. 2005). However, Cu readily forms precipitates on lattice defects which can make the gettering within a reasonable amount of time and temperature inefficient (Istratov and Weber 1998). Furthermore, it was reported that the recombination activity of Cu increases as Cu forms precipitates (Weber 1983, Baumann et al. 1997, Istratov and Weber 1998, Buonassisi et al. 2005). Indeed Cu tends to form precipitates or defect/impurity complexes in preference to existing at interstitial or substitutional sites in the Si lattice. As a result, Cu may be detrimental to Si solar cell efficiency even at low contamination levels of about 1012-1013 cm-3, which is lower than the tolerable limit reported in (Davis et al. 1980, Hopkins and Rohatgi 1986) (Figure 2.10).

Figure 2.10: Threshold impurity concentrations for CZ cell performance reduction (Istratov and Weber 1998) after (Davis et al. 1980, Hopkins and Rohatgi 1986).

29 In a recent study, Hernandez, et.al. investigated the effectiveness of different metal barrier layers deposited by physical vapour deposition (PVD) in preventing the diffusion of Cu in Si along with the impact of different thermal stresses on solar cell performance (Hernandez et al. 2010). Among other barrier metals (Ti, Ti/TiN, Ta and TaN) tested, a bilayer of Ni and Cu was deposited by PVD. Then, Cu electroplating was used to complete the metallisation process. The variation in the Voc of the samples after 1 minute of different thermal treatments was measured. One sample was used as a reference with Cu deposited directly on Si. The performance of the reference sample was degraded in temperatures below 100 °C. It was concluded that several hours of annealing at 230 °C, did not significantly degrade the performance of the solar cells. However, the silicidation process during metallisation can degrade the performance of the cells at high temperatures (Hernandez et al. 2010).

In another report, Bartsch et.al. studied the long term effects of a Cu-containing metallisation system on the performance of Si solar cells (Bartsch et al. 2010). Nickel barrier layers were deposited by electroless plating and light-induced electrolytic plating. The Suns-V oc technique was used to measure the pseudo fill factor (pFF) of solar cells which is highly dependent on the dark saturation current density Jo2. Since Jo2 is affected by the presence of impurities such as Cu in the space charge region (SCR), the variation of the pFF over a period of time (duration of the thermal stress) was used as an indication of the presence of Cu in the SCR. Therefore, it was assumed in this study that a certain loss in the pFF corresponds to a certain depth of Cu diffusion in the solar cell. It was shown that the pFF decreased markedly after several minutes of thermal stress at 300 °C. However, the pFF degraded at a much slower rate during a 200 °C thermal stress. It was concluded that the Ni barrier layer effectively prevents the diffusion of Cu into the solar cell, though disappointingly no data was given for how the thermal treatment affects the overall device efficiency (Bartsch et al. 2010).

30 2.7 Summary

In this chapter a review of previous studies which investigated the properties of Ni diffusion barrier layers for Cu, and the physical and electrical behaviour of Cu in Si was presented. Electrolessly deposited Ni-P alloys are frequently claimed to be the most effective Ni barriers layer in preventing Cu diffusion, owing to their amorphous structure. The formation of intermetallic phosphides in Ni barrier alloys is shown to be an important factor in retarding Cu diffusion. The dominant diffusivity paths in a crystalline material are grain boundaries and dislocations. Therefore, the co-deposited P from electroless Ni plating blocks the diffusion path since it precipitates uniformly along the grain boundaries within the alloy. The Ni-P alloy which is the main contributor to improving the barrier properties, is seemingly formed by annealing at 350 °C. It has been demonstrated that the properties of Ni-P alloys as diffusion barriers are improved as the P content is increased. An increase in the thickness of the Ni deposit also assists in blocking Cu penetration. The pure Ni deposited from electroplating is reported to restrict but not totally block Cu from entering the Ni barrier layer.

Studies on Cu behaviour in Si have revealed that Cu diffuses in Si as positively charged ions. Consequently, its diffusivity in Si depends on the conductivity and doping level of Si, as well as the temperature. It has been reported that after Cu has diffused into the bulk of a Si wafer during heat treatment, it may either precipitate in the bulk or out- diffuse to the wafer surface when the wafer is cooled down to room temperature. The Fermi level position in the Si wafer is shown to determine the preferred reaction path. Copper precipitates introduce band like states in the upper half of the Si bandgap and thus create efficient recombination sites. It was reported that Cu precipitation is more likely to occur in n-type Si than in p-type Si. Similarly, the recombination activity of Cu is stronger in n-type Si than in p-type Si. Furthermore, it was reported that Cu could be out-diffused from p-type Si wafers at room temperature if the surface oxide was removed while out-diffusion of Cu from n-type Si was reported to occur at 400 °C and above.

31 Precipitation of Cu in Si solar cells can degrade the electrical performance of the cells. Since Cu tends to precipitate on nucleation sites and defects in Si, multi-crystalline solar cells are probably less tolerable to Cu contamination than high quality CZ cells.

32 Chapter 3 Characterisation of Copper-rich Regions in Silicon Solar Cells

3.1 Introduction

The advent of several selective-emitter solar cell designs has allowed the use of an advanced self-aligning metallisation method based on plating of Ni and Cu. Despite its advantages over the screen-printed contacts, this method brings Cu in close contact with Si. Typically, the Si substrate is protected from the fast diffusing Cu by a reliable diffusion barrier layer (Frant 1961, Antler 1970 , Turn and Owen 1974). The use of Ni as a diffusion barrier for Cu and the impact of Cu contamination in microelectronic applications and semiconductor devices have been extensively studied as mentioned in Chapter 2. However, most of these studies focused on intrinsic Si, p-type or n-type Si wafers and p-n junction devices (Miyazaki 1991, Shabani et al. 1996, Istratov et al. 1998, Istratov and Weber 1998, Seibt et al. 1998, McHugo et al. 1999, Flink et al. 2000, Istratov et al. 2000a, Istratov et al. 2000b, Istratov and Weber 2002). While some attention has been given to solar cell devices (Buonassisi et al. 2005, Bartsch et al. 2010, Hernandez et al. 2010), few studies have investigated Cu diffusion through different types of Ni barrier layers. Furthermore, no studies have investigated the specific case of Cu diffusion through Ni barrier layers for LDSE solar cells.

This chapter reports on spectroscopy investigations into the effectiveness of Ni barrier layers, deposited by electroless plating and LIP, in preventing the penetration of Cu into the Si of LDSE solar cells. In the reported experiments, LDSE cells were heat-treated to intentionally diffuse Cu from the plated metal contacts through the Ni barrier layer and into the underlying Si wafer. The treated cells were then characterised using EDX, in conjunction with transmission electron microscope (TEM). The combination of EDX

33 and TEM was used as the primary characterisation method for chemical identification of Cu-rich clusters in the Si regions of the cells.

A brief review of the TEM and EDX analysis methods are presented in Section 3.1.1 and Section 3.1.2, respectively. In the remainder of the chapter a description of the conducted experiments is provided, followed by presentation and discussion of the TEM/EDX results.

3.1.1 Transmission Electron Microscopy (TEM)

Transmission electron microscopy enables very high resolution images to be obtained from ultra thin, electron transparent specimens. The high resolution imaging capability of TEM makes it a powerful technique for imaging and analysis of nanostructures and very fine particles. The TEM method makes use of an electron gun to create a beam of electrons to bombard a very thin sample. The electrons pass through the thin material section and the interaction of these electrons with the specimen provides information about the internal structure of the specimen. This information can be in the form of a conventional image or diffraction patterns.

The contrast in TEM images depends on the mode of operation. In bright field mode, the areas of the image which appear darker indicate the regions of higher atomic number or thicker regions of the sample, as fewer electrons are transmitted in these areas. If there is no sample in the beam path, the area will appear bright.

In TEM analysis, sample preparation can be a significant concern since accurate analysis requires good quality specimens. The TEM specimen preparation used for the experiments reported in this chapter is described in Section 3.2.1.2.

34 3.1.2 Energy Dispersive X-ray Spectrometry (EDX)

The energy dispersive X-ray spectrometry technique is used for identifying the elemental composition from a very small amount of material, or a very small part of a larger specimen. An EDX system requires a source of high-energy electrons for chemical microanalysis which can be provided by an electron microscope. Scanning electron microscopes (SEM) and TEMs are usually equipped with energy dispersive X- ray detectors to allow compositional analysis on such instruments. However, since TEM samples are very thin, EDX analysis in TEM offers a much higher spatial resolution than that in SEM. Figure 3.1 shows an electron beam hitting the surface of a specimen. For bulk specimens, the electrons become scattered in a lateral direction as they penetrate deeper in the specimen and form a pear-shaped “interaction volume” below the sample surface. The size of the interaction volume varies depending on the energy of incident electron beam and the atomic number of the element. Thus, it is possible to reduce the size of the interaction volume by reducing the energy of the electron beam. However, the energy of the incident electrons must be kept high enough so as to generate X-rays in the specimen. Therefore, the spatial resolution of X-ray microanalysis of bulk samples in SEM is limited to the minimum size of the interaction volume which is typically about 1 µm3 (Goodhew et al. 2001).

In TEM, however, the beam broadening is very weak. The TEM samples are very thin (~ 150-200 nm) and the electrons are transmitted through the specimen before they can spread much. Therefore, the spatial resolution in TEM is determined by the beam diameter which is ~ 5-10 nm. This allows for very small regions of ~ 10 nm or less in the specimen to be analysed.

An EDX system interfaced with TEM was used in the experiments reported in this chapter for qualitative elemental analysis.

35 Incident electrons

5-10 nm

Thin specimen ~100-200 nm Secondary electrons

Auger electrons

Backscattered electrons

Bulk specimen 1-2 µm Characteristic X-rays

Bremsstrahlung X-rays Fluorescent X-rays

Figure 3.1: The electron beam interaction diagram for bulk and thin specimens (Materials Evaluation and Engineering 2010).

The Sources of X-rays

Electron microscopes normally use a high-energy electron beam to excite the atoms of the specimen. By interaction of these electrons with the specimen atoms, a range of signals, namely secondary electrons, backscattered electrons, characteristic X-rays and bremsstrahlung are generated (Bruker-AXS 2006). A detailed description of secondary and backscattered electrons can be found elsewhere (Goodhew et al. 2001, Materials Evaluation and Engineering 2010).

An EDX analysis is based on the principle that the wavelength (energy) of a characteristic X-ray depends on the atomic number of the element from which they are emitted (Moseley‟s law). An energy dispersive X-ray detector is used to measure and analyse the energy of the X-rays which allows qualitative identification of the elements present in the specimen.

The intensity of the X-rays indicates the concentration of the elements within the sample (Goodhew et al. 2001). Therefore, measurement of the number of the X-rays emitted per second enables a quantitative analysis. However, quantitative measurements

36 and calculations of an element concentration from the recorded X-rays require a much more meticulous process and the results may not be accurate in absolute terms due to errors associated with the required mathematical procedure.

Characteristic X-rays

Characteristic X-rays are generated when the energy of incident electrons is high enough to knock out an electron from an inner atomic shell within the sample. The excitation energy of the incident electron (E0) must be greater than the critical energy of the relevant atomic shell (Ec) in order to ionize the atom and generate characteristic X- rays. The critical energy is a unique value for each atom and electron shell. The probability of ionization of a certain shell within an atom is maximum when the ratio of incident electron energy to the critical energy of the atomic shell, which is called the overvoltage ratio (U0= E0/Ec) is between 2 and 3 (Bruker-AXS 2006). As can be seen from Figure 3.2, the probability of ionization or ionization cross section slowly decreases at larger values.

Figure 3.2: Ionization cross section versus overvoltage ratio (Bruker-AXS 2006).

37 As the inner shell electron is knocked out a vacancy is created. Subsequently, an electron from an outer-shell falls into the vacant inner-shell position, releasing some of its energy by emitting an X-ray (Figure 3.3). The energy of these X-rays is determined by the difference in the excitation energies between the two shells involved in the electron transition and is dependent on the atomic structure of the element.

An EDX system typically produces an energy dispersive spectrum. This is a histogram of the number of X-rays detected versus the energy levels, which are typically expressed in keV. The position of peaks from the characteristic X-rays, which are displayed on the EDX spectrum, are unique for each atom and therefore reveal qualitative information about the composition of the sample. As mentioned earlier in this section, the intensity (height) of the peaks reflects the concentration of that element within the specimen.

a) b) Ejected electron Characteristic X-ray Incident electron beam Incident electron beam

K K

L L M Bremsstrahlung M

Figure 3‎ .3: Ionization and generation of characteristic and bremsstrahlung X-ray (from (Bruker-AXS 2006)).

Characteristic X-rays of various energies are typically evident in the EDX spectrum. The X-ray energy depends on which of the outer-shell electrons fills the inner-shell vacancies (Figure 3.4).

According to theory, the electrons of an atom are most likely to reside in distinct shells surrounding the nucleus. Each of these shells is characterised by discrete energy levels. The electron shells are labelled K, L, M, N, O, P, and Q; starting from the one closest to the nucleus and can be divided into one or more subshells (Figure 3.4). The electrons which are closest to the nucleus (e.g., K shell electrons) have lower energy as they are

38 strongly attracted by the nucleus and therefore a higher energy is required to displace them (Chalmers 1982).

If an electron is ejected from the K shell and is replaced with an electron from the L shell, a Kα X-ray is emitted; if it is filled by an electron from the M shell, a Kβ X-ray is emitted. Similarly, if a vacancy is created in the L shell and an electron from the M shell fills the vacancy, an Lα X-ray is generated; if the vacancy is filled by an electron from the N shell, an Lβ X-ray is generated.

As mentioned above, characteristic X-rays are generated when the incident electrons have energy in excess of the resulting X-ray energy. Thus, depending on the excitation energy of the electron beam, different X-ray lines can be used in EDX analysis. The Kα X-rays which are the strongest line in the K-series are mostly used in chemical microanalysis of lighter elements. However, since the energy required to knock out a K- shell electron increases with atomic number, it may be essential to use L-series X-rays to detect heavy elements.

Figure ‎3.4: Transitions leading to characteristic X-ray generation (Bruker-AXS 2006).

39 Bremsstrahlung

Bremsstrahlung (German for “braking radiation”), is emitted when the high-energy electrons are decelerated in the electric field of atomic nuclei and subsequently lose their energy in the form of an X-ray. In this case, the X-ray is no longer characteristic of a particular atom. As the amount of the energy loss of the electrons can vary from zero to the maximum electron beam energy, the energy distribution of the bremsstrahlung X- rays develops as a continuum spectrum background, rising towards lower energies. The upper limit of bremsstrahlung energy spectrum corresponds to the incident energy of the electron beam, i.e., when the incident electron loses all of its energy by hitting an atom and emitting a single X-ray energy. While bremsstrahlung X-rays contain no useful compositional information about the sample being investigated, they should be taken into account whenever a quantitative analysis is applied (Goodhew et al. 2001, Bruker- AXS 2006).

Auger Emission

As described earlier in this section, when an incident electron hits the atom of a specimen, an inner-shell electron is ejected leaving behind a vacancy. An electron from an outer-shell then fills the vacancy, resulting in a release of energy. Auger emission which is an alternative to X-ray emission may occur where the energy released by the electron is transferred to another electron in an outer shell causing it to eject from its atomic shell. The second ejected electron is called an Auger electron. Measurement of the energy of the Auger electrons is the basis of AES ( Goodhew et al. 2001).

Fluorescence

Characteristic X-rays may be generated by another mechanism other than direct excitation by incident electrons. A number of primary X-rays generated by electron excitation which are passing through a specimen may themselves excite atoms which can lead to the generation of lower energy X-rays. This is known as fluorescence effect and since it enhances the characteristic X-rays by some percent, it must be taken into account in quantitative analysis. The X-ray fluorescence is the primary signal source in X-ray fluorescence spectroscopy (XRF) (Goodhew et al. 2001, Bruker-AXS 2006).

40 3.2 Verification of the Characterisation Method

In order to verify if the sensitivity of the TEM analysis method is sufficient to detect Cu penetration in Si, three batches of two LDSE solar cells were fabricated as test samples. No diffusion barrier was used in the metallisation of these samples; instead, Cu was plated directly on Si. Section ‎3.2.1 describes the experimental procedure of making the LDSE cells and preparation of specimens for TEM analysis. The results from these experiments are presented and discussed in Section ‎3.2.2.

3.2.1 Experimental

3.2.1.1 Preparation of Laser-Doped Selective-Emitter Solar Cells

Commercial-grade p-type 1 Ω.cm CZ Si wafers were alkaline-textured, and then phosphorus-diffused to form an 85 Ω/□ n-type emitter. A 75 nm layer of Si nitride

(SiNx) was deposited using plasma-enhanced chemical vapour deposition (PECVD) as an ARC on the front surface. The approximate refractive index of the layer was 2.0. Aluminium was screen-printed on the rear surface of the wafer and then fired at a peak firing temperature of 860 °C for 3 seconds using an inline firing furnace to form a back surface field (BSF) and the rear metal contact.

The wafers were then uniformly spin-coated with 85% phosphoric acid (H3PO4) at speed rate of 3000 rpm for 40 seconds. The cells were then immediately laser-doped using a 532 nm high-powered laser with an average power of 15 Watts and scan speed of 5 m/s, to form selective-emitter regions. The laser patterns the dielectric layer while melting the underlying Si and simultaneously incorporating the dopants into the molten Si (Tjahjono et al. 2007a). The wafers were then rinsed in deionized (DI) water to remove any residual H3PO4 from the front surface.

Then, Cu was light-induced plated directly on the exposed Si in the laser-doped patterns. Figure 3.5 presents a schematic diagram of the LIP bath set-up used in all light-induced metal plating experiments in this Thesis. As described in Section 2.3, the light-generated current within the solar cell under illumination drives the LIP process. In this LIP process set-up, the Al rear of the wafers which is in contact with the p-type

41 surface of the wafers acts as the anode. Light-generated electrons are collected by the n- type emitter on the front surface of the wafers and are consumed by the reduction of metal (e.g., Ni and Cu) ions, present in the LIP solution, to form metal deposits. At the rear surface, electrons generated by the oxidation of Al, recombine with the light- generated holes to complete the circuit.

The anodic reaction for both the LIP Ni and Cu arrangements used in this Thesis is:

퐴푙 → 퐴푙3+ + 3 푒− 퐸0 = 1.66 푉

The cathodic reaction, that occurs in the laser-doped regions, for LIP Ni is:

푁𝑖2+ + 2 푒− → 푁𝑖 퐸0 = −0.25 푉 and for LIP Cu is:

퐶푢2+ + 2 푒− → 퐶푢 퐸0 = 0.34 푉

Light source n++ laser-doped surface (cathode) Electrolytic solution with metal salts Lightly diffused emitter Si solar cell

Al rear Supporting contact (p++) Teflon poles Figure 3.5: A schematic diagram of the LIP bath set-up used for all experimental work in this Thesis. A wafer with laser-doped regions is immersed in a petri dish filled with LIP solution of the metal to be plated, and exposed to constant light intensity from a CFL.

Prior to metallisation, the fabricated LDSE cells were first deglazed in 1% (w/v) hydrofluoric acid (HF) solution to remove any native oxide and phosphosilicate glass (PSG) from the exposed heavily-doped regions of Si and then rinsed in DI water. Then, the wafers were placed in petri dishes containing a Cu LIP solution, with the laser-

42 doped lines facing upwards and exposed to tube compact fluorescent light (CFL) sources. A proprietary Helios Cu formulation from MacDermid Inc1 was used for all Cu

LIP in this thesis. This formulation contained Cu sulphate (CuSO4.5H2O), sulphuric - acid (H2SO4), chloride ions (Cl ) and proprietary additives. The wafers were plated for 5 minutes at room temperature. Since the thick metal layer may lead to poor adhesion between the metal and the Si, particularly when there is no NiSi layer present, the shorter Cu plating time of 5 minutes was chosen for the test samples. After the plating, the cells were rinsed and dried, and were split into three batches.

The cells from the first and second batches were sintered in a muffle oven at 400 °C for 5 hours and 24 hours, respectively, and subsequently cooled under ambient laboratory conditions. The Cu in-diffusion temperature of 400 °C was chosen in order to diffuse sufficient concentrations of Cu above the detection limit of the TEM system. According to Equations 2.4 and 2.5 (Istratov et al. 2000a, Istratov and Weber 2002), the solubility of Cu in Si at 400 °C is as high as 3.8 × 1012 cm-3 and the diffusion coefficient of Cu in 1 Ω.cm p-type Si reaches 1.35 × 10-5 cm2/s. This diffusivity is high enough to enable Cu to traverse a 200 µm thick p-type Si wafer in less than 10 seconds.

The cells from the third batch underwent a heat treatment for 4 hours at 550 °C for Cu diffusion, followed by a quench in ethylene glycol. Since it was reported previously (Seibt and Graff 1988, Istratov and Weber 1998, Seibt et al. 1998) that the cooling rate of the heat-treated samples affects the size of Cu precipitates in Si, a fast cooling rate by quenching in ethylene glycol was examined in this experiment. Furthermore, due to the higher Cu solubility and diffusivity at this temperature, a shorter sintering time of 4 hours was chosen for these samples. Table 3.1 summarises the heat treatment parameters and cooling method used for the samples in this experiment. The bulk of the Cu contacts on all samples were removed prior to TEM investigation by etching in 50% nitric acid followed by full RCA cleaning. The TEM specimens were prepared right beneath the Cu lines from all samples as explained in Section 3.2.1.2‎ . The TEM/EDX studies of Cu diffusion and chemical identification results from these experiments are presented in Section ‎3.2.2.

1http://electronics.macdermid.com/products/photovoltaics/plated_metal/l 43 Table 3.1: Summary of heat treatment parameters and cooling method.

Heat Treatment Group Time Cooling Method Temperature (°C) (hr:mm) 1 400 5:00 Slow Cooled

2 400 24:00 Slow Cooled

3 550 4:00 Quenched

3.2.1.2 Specimen Preparation for TEM

Specimens for TEM analysis must be sufficiently thin to allow electrons to pass through the sample. A suitable TEM sample needs to have a thickness in the range of 20 to 200 nm. Apart from traditional techniques of sample preparation, single or dual focused ion beam (FIB) methods have been widely used for nanomilling very thin sections of material for TEM analysis. An FEI XP200 single beam FIB is used to prepare all TEM samples in this thesis. This system uses a focused beam of high-energy Gallium (Ga) ions to scan over the surface of a sample. To mill subsurface cross-sections from a specimen, the FIB system is operated at high currents. The high energy Ga beam strikes the sample and sputters away the atoms from the surface. At low beam currents very little material is removed by sputtering and FIB can be used for high resolution imaging. To prevent sample charging, the sample is firstly covered with a thin layer of ~25 nm of Au using an Au sputter coater prior to ion beam milling. Furthermore, in order to protect the underlying material from the destructive effect of high current beam during the milling process, a sacrificial layer of metal is deposited on all samples using the gas chemistry system interfaced with the FIB instrument. In all experiments in this work, a layer of ~1 µm platinum (Pt) was deposited on the sample where the thin section was to be milled.

Due to the nanometre scale resolution of the FIB, specimens can be precisely milled down to a scale of a few hundreds of nanometres. The TEM specimens prepared by the FEI XP200 single beam FIB in this work were approximately 18 µm by 5 µm and ~150 nm thick. Figure 3.6a shows a cross-sectional view of a TEM sample in the trench 44 milled by FIB. The last step in FIB is to detach the specimen from the bulk by cutting off the sides and bottom of the sample. The sample is then removed from the FIB-cut trench and placed onto a standard TEM supporting grid. Due to the impact of the X-rays generated from the supporting grids on the measured X-ray spectrum, an Au grid is used in this work, rather than a Cu grid. The TEM samples were lifted out using a micro- manipulator under a light microscope. This method was difficult to perform and frequently resulted in lost samples (i.e., poor yield). Figure 3.6b shows an optical microscope image of a TEM sample mounted onto an Au grid.

a) b)

Figure ‎3.6: a) A FIB image of a TEM specimen prepared by a single beam FIB system, b) corresponding optical microscope image from the TEM specimen mounted on a supporting Au grid.

Although FIB can be used to nanofabricate TEM samples precisely, the milling process may introduce damage on the side walls of the specimen. This effect is slightly alleviated by the use of a low current beam on the final stage of milling. Unfortunately, samples prepared by FIB tend to have a non-uniform thickness. Furthermore, the time required for sample preparation in the single beam FEI XP 200 is almost double the time required as in a dual beam FIB.

All TEM samples for this thesis were prepared by the author using the above mentioned techniques.

45 3.2.2 Results and Discussion

When no Ni barrier layer was present, the TEM/EDX analysis on the first batch of LDSE solar cells which were heat-treated at 400 °C for 5 hours and subsequently cooled under ambient laboratory conditions, did not reveal Cu penetration into the Si substrate. This observation may be due to the slow cooling conditions and the short sintering time which together may have resulted in Cu concentrations in Si below the detection limit of EDX. The influence of cooling conditions on the process of Cu precipitation is discussed in details in Section 4.5.2.

On the other hand, the TEM images of the second and third batches of heat-treated LDSE cells with no Ni barrier layer, revealed dark regions close to the surface of the Si in both groups. As described in Section 3.2.1.1‎ , the second batch was heat-treated for 24 hours at 400 °C and subsequently cooled in air whereas the third batch underwent heat treatment for 4 hours at 550 °C followed by a quench in ethylene glycol. The TEM images presented in Figure ‎3.7 clearly show the relatively large dark features with diameters in the range of ~ 50 to 200 nm in a sample from the second group. The observed dark areas were randomly distributed close to the surface of Si, extending up to ~ 500 nm deep and were determined as Cu-rich regions by EDX analysis. Several analysis methods such as individual point analysis, line scan and elemental mapping were performed to obtain and evaluate energy dispersive spectra from the defined areas.

46 a) b)

Pt 200 nm

c) d)

Figure 3.7: TEM images of Cu/Si film from a sample in the second group with no Ni barrier layer, after heat treatment at 400 °C for 24 hours and cooled in air, a) exhibiting dark features which were verified as Cu-rich regions by EDX analysis, b) a close-up view showing the large diameter of the Cu particles, c) the Cu particles extending to ~ 350 nm from the Si surface, and d) a close-up view the Cu particles.

In the elemental mapping method, an electron image of the area of interest is first captured (Figure 3.8a). While the electron beam scans the surface of the area continuously, the corresponding X-ray data are collected point-by-point and simultaneously recorded into the map database. This allows individual colour-coded elemental maps to be obtained which displays the concentration distribution of predefined chemical elements across the area of the specimen. Figure 3.8b and Figure 3‎ .8c present the elemental maps for Cu from the specific area of a sample from the second group. The average energy dispersive spectrum from the whole mapped area is also shown in Figure ‎3.8d. As can be seen, the peak X-ray energy positions of Cu-Kα1

47 at 8,047.78 eV and Cu-Lα1 at 929.7 eV (Thompson et al. 2009) together with the elemental maps verify the presence of Cu in those regions.

Small peaks from iron (Fe) of Kα1 at 6,403.84 eV and Kα2 at 6,390.84 eV and cobalt

(Co) of Kα1 at 6,930.32 eV and Kα2 at 6,915.30 eV (Thompson et al. 2009), in the EDX spectrum, are artefacts originated from the TEM pole piece and are common in EDX analysis. Carbon (C) and oxygen (O) are also commonly observed in EDX spectra.

Gallium peaks of Kα1 at 9,251.74 eV and Kα2 at 9,224.82 eV (Thompson et al. 2009), originate from FIB ion milling. Gold peaks of Lα1 at 9,713.3 eV and Lα2 at 9,628.0 eV (Thompson et al. 2009) were collected from the supporting Au grid.

48 a) b)

BF Cu

Figure 3‎ .8: a) An electron image from the area of the specimen with a dark spot, from a sample in the second group with no Ni barrier layer, after heat treatment at 400 °C for 24 hours and cooled in air, b) and c) corresponding X-ray elemental maps from the defined area, d) average energy dispersive spectrum obtained for the whole scan area.

49 Figure ‎3.9a shows a TEM image of a sample in the third group after sintering at 550 °C for 4 hours followed by a quench in ethylene glycol. Using EDX, Cu was found on the dark features which extended up to ~ 750 nm deep from the surface of the Si. Figure ‎3.9b is an example energy dispersive spectrum obtained using EDX individual point analysis. It clearly shows the visible peaks of characteristic X-ray energies of Cu-Kα1 at

8,047.78 eV and Cu-Lα1 at 929.7 eV (Thompson et al. 2009).

In order to confirm the hypothesis that the Cu peaks observed in the EDX spectra originated from the plated Cu contacts diffusing into the Si due to the heat treatment, a few individual point analyses were performed on different regions of the samples. Figure ‎3.10a presents the electron image of a sample that was subjected to EDX point analysis. EDX was performed at the marked points. Figure 3.10b is an example of the energy dispersive spectrum of each point which revealed no reliable Cu peaks. The results from this analysis, and the fact that no Cu peaks were detected in the first group that were only heated for 5 hours suggest that artefacts from the microscope or the TEM/EDX process did not result in false Cu peaks. Furthermore, the background concentration of Cu in CZ Si is typically in the order of 1013 cm-3 (Bai et al. 1990), which is less than the detection limit of 0.1 at.% typical of electron probe microanalysis techniques (Gauglitz and Vo-Dinh 2003). Therefore, the results of this experiment indicate that the Cu detected by the EDX analysis method is not affected by the background Cu concentration in the wafers. As a result, it was concluded that the Cu detected in the samples in groups 2 and 3 had diffused into the Si from the Cu metallisation regions. However, clearly it was not possible to test all points in each image so the failure to detect Cu in the group 1 samples cannot be used to conclude that Cu did not diffuse into the Si in those cells.

50 a)

Figure ‎3.9: a) A TEM image of a sample in the third group with no Ni barrier layer, after heat treatment at 550 °C for 4 hours followed by fast quenching, b) energy dispersive spectrum obtained from an individual point analysis on the dark features.

51 a)

Figure 3.10: a) An electron image from a sample with no Ni barrier layer, after heat treatment at 550 °C for 4 hours followed by fast quenching. Points subjected to EDX analysis are marked, b) an example EDX spectrum obtained from individual point analyses on the marked areas.

It is important to note that some of the Cu-rich clusters observed in these samples may have been new precipitates grown on the defects in the Si created by ion milling. As mentioned in Section ‎3.2.1.2, the ion milling process by FIB to prepare the TEM specimens may introduce structural defects or deformities on the side walls of the specimen. As Cu remains mobile at room temperature (Weber 1983, Istratov et al. 1998, Flink et al. 2000, Istratov et al. 2000a, Istratov and Weber 2002), it is possible that Cu may have agglomerated in the defects or deformities and formed precipitates, during or 52 after the milling process, before starting TEM/EDX investigations. Nevertheless, irrespective of whether the defects are induced during milling, heat treatment or previously existed in the sample, the formation of precipitates in the Si demonstrates that Cu has indeed diffused into the Si in groups 2 and 3 and can be detected in the form of precipitates by EDX.

3.2.3 Conclusion

It was concluded from the TEM/EDX investigations that when no Ni barrier layer was present, the 5 hour heat treatment at 400 °C in combination with subsequent slow cooling in air, did not result in a Cu concentration in Si that could be reliably detected by TEM/EDX. On the other hand, the TEM/EDX investigations revealed that Cu diffuses into the underlying Si during a long heat treatment of 24 hours at 400 °C followed by a slow cool or a 5 hour heat treatment at 550 °C followed by quenching in ethylene glycol. More importantly, it was established that TEM/EDX analysis can successfully detect the diffused Cu at this level of contamination. Therefore, although this technique was not used quantitatively to determine the concentration of Cu in Si, it is well suited for qualitatively detecting the diffused Cu due to Ni barrier failure in LDSE solar cells.

On the basis of these experimental results, a maximum sintering temperature of 400 °C was chosen to perform experiments to detect the failure of LIP Ni and electroless Ni barrier layers in LDSE cells (Section ‎3.3.1.1 and Section ‎3.3.1.2).

Also, according to the previous reports (Seibt and Graff 1988, Istratov et al. 1998, Istratov and Weber 1998, Seibt et al. 1998), the cooling rate of the heat-treated samples affects the size of Cu precipitates in Si. Also, at high cooling rates of greater than 100 K/s, Cu precipitates uniformly in the bulk of the wafer as Cu does not have time to diffuse to the Si surface, whereas slow cooling rates (e.g., cooling in air), result in Cu precipitation in regions close to the wafer surface. Based on the observations from this section and the abovementioned previous studies, a fast cooling procedure by quenching in ethylene glycol was chosen to investigate and compare the diffusion of Cu through different types of Ni barrier layers. Istratov, et.al, had estimated the cooling rate of

53 around 1000 K/s for ethylene glycol (Istratov et al. 1998). The corresponding experimental work is described in Section ‎3.3.1.1 and Section ‎3.3.1.2.

3.3 Comparison between Electroless Nickel and Light-Induced Plated Nickel as Diffusion Barrier Layers

3.3.1 Experimental

Two groups of LDSE cells were fabricated using the wafers prepared as described in Section 3.2.1.1 but with differing metallisation steps as shown in Table ‎3.2. The front metal contacts were formed by a self-aligned metallisation scheme of Ni plating, Ni sintering and Cu plating. The individual metallisation steps are described in more details below.

Table 3.2: Summary of metallisation groups, heat treatment parameters and cooling method.

Heat Treatment Cooling Group Ni Barrier Layer Cu Layer Temperature (°C) Time (hr:mm) Method

400 5:00 Quenched 1 LIP LIP Helios Cu 200 5:00 Quenched

400 5:00 Quenched 2 Electroless LIP Helios Cu 200 5:00 Quenched

3.3.1.1 Metallisation Process of Light-Induced Plating

Laser-doped selective-emitter cells were fabricated as described in Section 3.2.1.1. Then, LIP Ni was performed to deposit a thin layer of Ni of less than 1 µm on the laser- doped lines. After the deglazing step, the wafers were immersed in the Ni plating solution, substantially as described for the LIP Cu step described in Section 3.2.1.1‎ , under a constant illumination from a CFL source. A Barrett SN1 Ni sulphamate solution supplied by MacDermid Inc. was used for all light-induced Ni plating in this thesis. The

54 Ni sulphamate bath comprised Ni sulphamate (Ni (SO3NH2)2), Ni chloride

(NiCl2.6H2O), boric acid (H3BO3) and proprietary additives. The plating bath temperature was kept at 40 °C.

The thickness and homogeneity of Ni deposits are important for metal plated contacts. A thick Ni deposit reduces the adhesion between the plated metal contact and the Si wafer because of the increased stress of the plated metal. In addition, the Ni deposit needs to homogenously cover the laser-doped regions on the surface of the Si wafers in order to inhibit Cu from diffusing into the underlying Si. To investigate the uniformity and thickness of the light-induced plated Ni, three LDSE wafers were plated with Ni via LIP with differing plating time of 1 minute, 40 seconds and 30 seconds. The thickness and the uniformity of the Ni deposit were examined using single beam FIB.

The 1 minute plating time resulted in a relatively thick Ni deposit of ~ 1.5 µm. Also, as can be seen in Figure 3.11a, Ni did not plate uniformly on the laser-doped lines. Similarly, the Ni layer from the 30 seconds of LIP did not cover the whole surface of the laser-doped lines. Figure 3.11b clearly shows large areas on the laser-doped region where Ni had not been plated. A Ni plating time of 40 seconds resulted in a homogenous Ni deposit with a suitable thickness of ~ 1 µm (Figure 3.11c and Figure 3.11d).

55 a) b)

c) d)

~ 1 µm

Figure 3.11: Top view FIB images of light-induced plated Ni for a) 1 minute, b) 30 seconds, c) close-up view of light-induced plated Ni for 40 seconds, tilted at 44.6 degrees, d) top view of Ni LIP for 40 seconds.

Based on the above observations, a plating time of 40 seconds for Ni LIP was chosen for all experiments. After Ni plating via LIP, the wafers were rinsed and dried thoroughly with a N2 gun prior to Ni sintering which was carried out to form a NiSi layer on the laser-doped lines. The formation of a mono NiSi layer during the sintering step is important as it reduces the contact resistance between the metal and Si interface, improves adhesion between the Si and the metal contact and prevents Cu from diffusing into the underlying Si (Wenham 1986, Deng et al. 1997, Gambino and Colgan 1998, Kim 2005, Fisher 2007). The wafers were sintered in a conventional tube furnace at 350 °C for 3 minutes in N2 environment.

56 Finally, the wafers were plated with LIP Cu for 10 minutes at room temperature and under a constant illumination from a CFL source, using the Helios Cu formulation from MacDermid Inc2. The resulting Cu fingers were about 10 µm thick. The cells were rinsed and dried and then were split into two groups. The cells from each group were heat-treated in a muffle oven for 5 hours at varying temperatures of 400 °C and 200 °C, respectively. All heat-treated samples were quenched in ethylene glycol. Table 3.2 presents a summary of heat treatment parameters and cooling method.

After heat treatment, TEM specimens were prepared from the Si surface immediately below the front metal contacts from all samples, using an FIB system as described in Section‎3.2.1.2. It was essential to remove the thick layer of the front metal contacts beforehand by chemical etching as the depth of the TEM specimens prepared by FIB milling is typically limited to ~ 4 to 5 µm. Data was collected using a Philips CM200 TEM system operating at 200 kV acceleration voltage. The EDX analysis was conducted using a Bruker-AXS UHV detector with a Si(Li) crystal and energy resolution of 129 eV, measured on full width of peaks from manganese, Mn-Kα, at half maximum (FWHM).

In order to perform X-ray microanalysis using a TEM, the sample should be tilted against the horizontal plane towards the EDX detector. A tilt angle of ~ 15 degrees is used for all EDX measurements on TEM in this work.

Results from TEM/EDX studies of Cu diffusion and chemical identification on these experiments are presented and discussed in Section ‎3.3.2.1.

3.3.1.2 Metallisation Process of Electroless Plating

To investigate the effectiveness of an electroless Ni barrier layer in preventing the diffusion of Cu into Si in LDSE solar cells, LDSE wafers were fabricated as described in Section 3.2.1.1. After deglazing, a thin layer of Ni of less than 1 µm was deposited on the laser-doped regions of the LDSE wafers by electroless Ni plating, using a similar process to that described in Section 2.2.

2http://electronics.macdermid.com/products/photovoltaics/plated_metal/l 57 Two LDSE wafers were used to examine the uniformity and thickness of the electroless Ni deposit with differing plating time. The wafers were immersed in an electroless Ni plating bath of 85 °C for 5 and 8 minutes. An ammonia type electroless Ni plating solution supplied from Transene Company, Inc.3 was used in all electroless Ni plating experiments in this work. The electroless Ni deposits were checked under SEM.

Figure 3.12: Top view SEM images of electroless Ni plating for a) 5 minutes, b) 8 minutes.

As can be seen from Figure 3.12, the electroless Ni layer after 8 minutes of electroless plating is more uniform than the one after 5 minutes of plating, with almost no visible holes on the Ni deposit. Also, the thickness of the Ni layer in this case was slightly less than 1 µm which is desirable in the metallisation of solar cells. Therefore, 8 minutes of electroless Ni plating was chosen for the LDSE cells in this group.

After electroless Ni plating, the wafers were rinsed and dried thoroughly. Nickel sintering was carried out in a conventional tube furnace at 350 °C for 3 minutes in N2 ambient. The sintering process was immediately followed by light-induced Cu plating to deposit a thick layer of Cu which acts as the bulk of the front metal contact.

The subsequent steps in the metallisation process and the parameters for heat treatments are described in Section 3.3.1.1. Following the heat treatments, TEM specimens were prepared as described in Section 3.2.1.2. Results from TEM/EDX investigations are presented and discussed in Section 3.3.2.2.

3http://www.transene.com/ni_ammonia.html 58 3.3.2 Results and Discussion

3.3.2.1 Analysis of Copper Diffusion through Light-Induced Plated Nickel Barrier Layer

The EDX/TEM characterisation of heat-treated LDSE solar cells with light-induced plated Ni as a diffusion barrier was performed on the samples sintered at 400 °C and 200 °C. In both cases it was found that Cu had diffused through the Ni layer into the underlying Si. Figure 3.13 and Figure 3.16 show TEM images of the sample heat- treated at 400 °C for 5 hours, followed by a quench in ethylene glycol. Copper was detected from EDX analysis in the small dark spots as visible in the TEM images. As can be seen from Figure 3.16a, these Cu-rich regions consisted of closely packed small particles and protruded up to ~ 900 nm deep from the surface of the Si. The diameter of the small Cu precipitates varied within the range of ~ 30-50 nm, as apparent from Figure ‎3.13b and Figure 3.16b.

a) b)

Figure 3.13: TEM images of a sample with LIP Ni barrier layer, a) Cu precipitates in the sample after heat treatment at 400 °C for 5 hours followed by a rapid quench, b) close-up view of a small sized Cu precipitate.

59 Various EDX point analysis and elemental mapping were conducted on the dark particles to verify their elemental composition. Figure 3.14, Figure 3.16c and Figure 3.16d are examples from the several EDX point analyses carried out on the heat- treated sample at 400 °C for 5 hours. The results from the point analysis indicate that these dark spots are mainly composed of Cu. The elemental mapping as presented in Figure 3.15, further demonstrates the presence of Cu in those regions.

Figure 3.14: a) An electron image from the heat-treated sample with LIP Ni barrier layer at 400 °C for 5 hours and fast quenching, showing an individual point on the closely-packed dark spots on the surface of the specimen, b) corresponding EDX spectrum obtained from the point analysis.

60 a) b)

c) d)

Cu BF Si Cu

Figure 3.15: a) An electron image from the small Cu precipitates on the specimen surface, from a sample with LIP Ni barrier layer, after heat treatment at 400 °C for 5 hours and fast quenching, b) and c) corresponding colour-coded X-ray elemental maps from the area, d) a multi-colour image formed from the superimposed elemental maps.

61 a) b)

Figure 3.16: TEM images of a sample with LIP Ni barrier layer after heat treatment at 400 °C for 5 hours with rapid quenching, a) showing the closely-packed small Cu precipitates reaching up to ~ 900 nm deep from the surface of the Si, b) close-up view of the small sized Cu precipitates, c) an electron image showing an individual point on the dark spots on the surface of the specimen, d) EDX spectra obtained from the individual point analysis.

62 The Cu-rich clusters detected by TEM/EDX analysis were located close to the surface of the Si. The depth of these regions in most cases did not exceed 1.8 µm. This suggests that the Cu precipitates were formed in the localized heavily-doped regions (n++ area) of the LDSE cell since the junction depth below the laser-doped lines was about 3-4 µm. The junction depth of an LDSE cell appears as a highlighted area in the Electron Beam induced Current (EBIC) image shown in Figure ‎3.17. It is also possible that some Cu that was trapped in the p-type side of the cell was released and diffused to the n-type Si where stable Cu precipitates form. Since the LDSE cells were fabricated from single- crystalline Si wafers with very few defects, Cu would more likely diffuse towards regions with more stable sinks. This could occur after the samples were cooled down to room temperature. These observations are in agreement with the previous findings that the formation of stable Cu precipitates is facilitated in n-type Si by the Fermi level position and possibly through formation of stable chemical complexes of Cu and P in n- type Si (Shabani et al. 1996, Istratov and Weber 2002). The precipitation of Cu in the emitter can severely reduce the minority carrier lifetime by providing effective recombination sites in solar cells (Baumann et al. 1997, Istratov et al. 1998, Istratov and Weber 1998, Seibt et al. 1998, Istratov et al. 2000b).

In addition, there is a possibility of Cu precipitation occurring in the lightly-doped emitter region or along the junction area of the LDSE cells provided the Cu diffused laterally. This is because the junction depth of the lightly-doped emitter is only ~ 0.5 µm, as apparent from Figure ‎3.17. This can deteriorate the p-n junction or create shunts within the solar cells.

On the other hand, due to the limitation of the small size TEM specimen which is typically up to 5 µm deep, it was not possible to inspect deeper in the bulk of the Si of LDSE cells. Therefore, if the Cu was gettered in the p+ area of the cell as explained in (Istratov et al. 2000a, Istratov and Weber 2002), or gettered at the Al back contact as reported in (Buonassisi et al. 2005), this detection was not possible using the TEM sample preparation method employed in this work.

63 ~ 0.5 µm ~ 3-4 µm

Figure 3.17: SEM-EBIC image of the cross-section of a LDSE cell revealing the junction depth below the laser melted region.

Although the main focus of this study was not to determine the chemical state of Cu- rich clusters, according to the previous findings (Nes and Lunde 1972, Das 1973, Nes and Washburn 1973, Solberg 1978, Seibt and Graff 1988 , Rizk et al. 1994, Istratov and

Weber 2002, Buonassisi et al. 2005), a form of Cu-rich silicide (Cu3Si) is the most dominant composition of the Cu precipitates in Si.

The TEM/EDX investigations here represent strong evidence for Cu diffusion in heat- treated LDSE solar cells through a LIP Ni barrier layer. Copper-rich precipitates distributed homogenously in the bulk of Si were visible in TEM observations even when samples were heat-treated at lower temperatures. From TEM images shown in Figure 3.18, it is evident that Cu in-diffusion at 200 °C for 5 hours with subsequent rapid quenching, resulted in formation of Cu precipitates as deep as 1.8 µm (Figure 3.18c) from the surface of the Si which is within the n++ region of the LDSE cell. The observed dark areas were verified as Cu-rich regions by performing several EDX analyses.

64 a) Pt b) Si

10 nm

c) d) Pt Si

20 nm Figure 3.18: TEM images of Cu particles in the heat-treated sample with LIP Ni barrier layer at 200 °C for 5 hours followed by fast quenching.

While it is possible that Cu can diffuse through the Ni barrier layer, another possible reason for Cu penetration may be defects in the Ni barrier layer which can be due to non-uniform Ni plating. As mentioned previously, the diffusion coefficient of Cu in 1 Ω.cm p-type Si at 200 °C is approximately 3.5× 10-6 cm2/s. Copper can diffuse a large distance of 5 mm in p-type Si wafer within 5 hours of sintering at this temperature. Furthermore, Cu remains mobile at room temperature (Weber 1983, Flink et al. 2000, Istratov et al. 2000b, Istratov and Weber 2002). This means Cu can penetrate into the underlying Si from a local penetration spot on the non-uniform Ni barrier layer, where Ni has not been plated. Therefore, the uniformity of the diffusion barrier layer may be critical in solar cell metallisation when Cu metal plating is employed (Istratov et al. 2000b, Istratov and Weber 2002, Tjahjono et al. 2010).

65 Another plausible explanation for the penetration of Cu into the Si substrate may be associated with possible defect spots on the ARC layer caused by the laser (Hernandez et al. 2010, Tjahjono et al. 2010). During the laser doping process, the ARC layer around the edges of the laser-doped area may be damaged or unintentionally removed. This can introduce defects on the ARC layer in the form of holes or cracks (Sugianto et al. 2007). As a result, the underlying lightly-doped emitter may be directly exposed to the plated Cu (Figure 3.19).

1) Dielectric 2) Holes in ARC peels off dielectric

n+ Laser-doped region

Figure 3.19: A cross-sectional schematic diagram of the exposed lightly-doped emitter on the perimeter of the laser-doped line (Used with the permission of Tjahjono 2010).

3.3.2.2 Analysis of Copper Diffusion through Electroless Nickel Barrier Layer

The EDX/TEM investigations of heat-treated LDSE solar cells with electroless Ni as a diffusion barrier demonstrate that Cu can also diffuse through these Ni barrier layers into the bulk. Figure 3.20a is a TEM image from the heat-treated sample at 400 °C for 5 hours with a subsequent quench in ethylene glycol, showing a large cluster, around 200 x 500 nm in size, containing closely packed dark regions and reaching up to ~ 790 nm deep in the Si bulk. The results from the EDX point analysis performed on these particles are presented in Figure 3.21 and prove the presence of Cu in these regions. Figure 3.21c and

Figure 3.21d clearly show both Cu peaks of Kα1 at 8,047.78 and Lα1 at 929.7 eV.

The Cu-rich clusters observed in this sample exhibited similar morphology and spatial distribution of Cu precipitates in Si, to the ones described in Section 3.3.2.1 for samples with LIP Ni as diffusion barriers. The Cu-rich clusters in both groups extended up to ~ 66 900 nm deep in the bulk of the Si. However, as can be seen from Figure 3.20d, the particles here appear to be smaller compared to those seen in Figure 3.13b and Figure 3.16b, with diameters in the range between 5 to 10 nm.

a) b)

c) d)

Figure 3.20: TEM images of the heat-treated sample with electroless Ni barrier layer at 400 °C for 5 hours and a fast quench, showing a) the closely-packed small Cu precipitates extended ~ 790 nm deep from the surface of Si, b) and c) close-up view of the Cu precipitates, d) a small Cu precipitate with 10 nm diameter.

67 a)

200 nm

b) cps/eV 3.5 Cell-Z

3.0

2.5

Ga 2.0 Co Fe Cu Si Fe Co Cu Ga

1.5

1.0

0.5

0.0 2 4 6 8 10 12 14 16 18 20 keV

c) x 0.001 cps/eV Cell-Z

500

400

Co Ga Fe Cu Si 300 Ca Fe Co Cu Ga Ca

200

100

0 4 5 6 7 8 9 10 d) keV

x 0.001 cps/eV Cell-Z

500

400

Fe 300 O Co Cu Ga Si Ca Fe Co Cu Ga Ca

200

100

0 0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 keV Figure 3.21: An electron image showing an individual point on the dark region on the heat-treated sample with electroless Ni barrier layer at 400 °C for 5 hours with a rapid quench, b) corresponding energy dispersive spectrum obtained from the individual point analysis, c) visible Cu peak of Kα1 at 8,047.78 eV and d) Lα1 at 929.7 eV on the zoomed spectrum x and y axes.

68 Further EDX/TEM studies revealed that sintering of LDSE samples at a temperature of 200 °C also led to a penetration of Cu in Si. Copper precipitates were observed to be uniformly distributed close to the surface of the sample. Figure 3.22 demonstrates the spatial distribution of the Cu precipitates in the heat-treated sample at 200 °C for 5 hours followed by quenching in ethylene glycol, along with the corresponding EDX analysis. Also visible in Figure 3.22d is the bremsstrahlung background which is shown in grey on the zoomed spectrum of the x and y axes. As can be seen, the peak of Cu-Kα1 at 8,047.78 eV is well above the background.

a) b) Pt

c) d)

Figure 3.22: a) A TEM image of Cu precipitates in the heat-treated sample with electroless Ni barrier layer, at 200 °C for 5 hours, with fast quenching, extending ~ 450 nm from the Si surface, b) an electron image showing an individual point on the dark area, c) corresponding energy dispersive spectrum obtained from the individual point analysis, d) zoomed spectrum in x and y axes exhibiting the bremsstrahlung background.

69 3.4 EDX Analysis Artefacts and Common Problems

When performing an EDX analysis, a number of common problems or artefacts need to be considered. These possible problems are described below and their potential contribution to erroneous analysis was assessed.

 Escape peaks: When trying to identify very small peaks, it is critical not to confuse the small peaks with escape peaks. The generated X-rays are measured by an energy dispersive X-ray detector. The detector used in the EDX analysis system in this work was Si(Li) which uses a lithium-drifted Si crystal, as the X-ray sensitive element. There is a possibility that the incoming X-rays excite the Si crystal in the detector and generate a Si-Kα X-ray of 1.74 keV. These X-rays could escape from the detector; therefore subtracting themselves from the full energy of measured X-rays of detected elements. This means that small peaks of energy 1.74 keV below the major characteristic X-ray peaks are expected to be observed in the EDX spectrum (Bruker-AXS 2006). These energy peaks are known as escape peaks and are typically about 0.5-1% of the major peak intensity.

The Au peaks observed in most EDX spectrum in this work, were collected from

the Au supporting grid. The main L line X-ray transitions for Au are Lα1 at

9.713 keV and Lα2 at 9.628 keV, which place the escape peaks at 1.74 keV below this energy range to about 7.973 and 7.888 keV, close to where the Cu-

Kα1 appears at 8.047 keV.

This artefact was avoided in the EDX analysis in this work as the Quantx system removes the escape peaks by automatic spectrum correction (Bruker-AXS 2006).

 Pile-up peaks: Pile-up or sum peaks are generated at higher count rates. If two X-rays hit the detector crystal simultaneously or very close in time, the measured X-ray energy

70 will be the sum of the two X-rays. The pile-up rejector in the spectrometry signal processor can distinguish and separate the piling-up of X-rays, if they come close in time. However, if the X-rays come too close in time or simultaneously, the pile-up rejector may not be able to suppress the pile-up peaks. Spurious pile-up peaks can be minimized by lowering the count rate (Bruker-AXS 2006). The pile-up peaks were avoided in the EDX analysis in this work by performing the measurements at low count rates.

 Sample drift: Sample drifting in all EDX analyses were minimal. In addition, the electron image may drift while performing EDX analysis which can affect the accuracy of the results. The drift correction feature in the EDX software used (Espirit) corrects for a small amounts of drift in the sample.

 Oven contamination test: As explained in Section ‎3.2.2, small peaks of Fe and Co observed in most EDX spectra are artefacts originating from the microscope polepiece. Similarly, the Au peaks were collected from the Au supporting grid used in all experiments. This hypothesis is further supported by an inspection of the contamination probability of the muffle oven which was used for heat treating the samples in all experiments described in this thesis. Particles from the surface of the base and walls inside the oven were collected and examined under SEM/EDX. Figure 3.23 shows the EDX spectrum obtained from several point analysis on the particles. As can be seen, there are no traces of the abovementioned elements which confirms the origin of these elements from the instrumentation parts of the TEM system.

71 a) b) Point-1

c) d) Point-2 Point-3

Figure 3.23: a) An SEM image of small particles collected from the muffle oven, b) EDX spectrum of point 1, c) EDX spectrum of point 2, and d) EDX spectrum of point 3.

3.5 Conclusion

In this chapter, the effectiveness of Ni barrier layers deposited by electroless plating and LIP, in blocking the diffusion of Cu in the Si bulk of LDSE solar cells was investigated. An EDX system in conjunction with a TEM was employed for chemical identification of Cu–rich particles in the Si of heat-treated LDSE cells.

It is shown that during the heat treatment at 400 °C or 200 °C (5 hours, followed by quenching in ethylene glycol), Cu penetrates into the bulk Si of LDSE cells through both electroless Ni and LIP Ni barrier layers. The TEM observations discussed in Section 3.3.2.1 and Section 3.3.2.2 reveal similar results in the morphology and spatial distribution of Cu precipitates in Si, for samples with LIP Ni and electroless Ni as diffusion barriers; although the size of the precipitates is smaller in the electroless- plated cells (~ 10 nm) than in the LIP cells (~ 30-50 nm). The Cu-rich clusters were in the form of closely-packed small precipitates, in regions close to the surface of Si. The presence of Cu in these regions is evidenced by several EDX analyses performed on LDSE samples. According to the literature, Cu precipitation is favoured in n-type Si by the Fermi level position (Istratov et al. 1998, Flink et al. 2000, Istratov et al. 2000a).

72 Also, the Cu precipitates are more stable in n-type Si probably through formation of stable chemical complexes of Cu and P (Shabani et al. 1996, Istratov and Weber 2002). This is in good agreement with the experimental data of this chapter. In most cases, the Cu precipitates extended up to ~ 1.8 µm deep from the surface of the wafers, which is within the heavily-doped selective-emitter region under the laser-doped lines of the LDSE cells. Most probably, the chemical state of the Cu precipitates is in the form of a

Cu-rich silicide (Cu3Si), which introduces band like states in the middle of the Si bandgap. Therefore, Cu precipitates can severely reduce the minority carrier lifetime by providing effective recombination sites (Istratov et al. 1998, Flink et al. 2000, Istratov et al. 2000a). However, the concentration of the minority carriers in the heavily-doped regions (the n++ area) in the LDSE structure is relatively low. Consequently, the formation of Cu precipitates in these regions may not cause significant problems. On the other hand, if Cu diffuses laterally, it may form precipitates in the lightly-doped emitter layer (the n+ area) or along the junction area. This can shunt the LDSE solar cells or cause the p-n junction of the cells to deteriorate. Thus, the formation of Cu precipitates within the lightly-doped emitter is likely to be detrimental to the electrical performance of solar cells. Due to the textured surface of the cell structure used in the experimental work of this thesis and the limitation of the TEM sample preparation equipment, it was not possible to characterise the lightly-doped emitter layer by the EDX/TEM method. However, it is recommended in the future to examine the lateral diffusion of Cu in future studies by performing similar EDX/TEM analyses in regions adjacent to the laser-doped regions.

Apart from the possibility of the diffusion of Cu through the Ni barrier layers, Cu penetration in the Si of LDSE solar cells may also occur as a result of non-uniform Ni barrier layers deposited by either electroless plating or LIP. A gap in the barrier layer, due to non-uniform plating, exposes the Si to the overlying Cu. Since Cu is mobile even at room temperature, such gaps may act as diffusion paths for Cu. Similarly, defects in the ARC layer around the edges of the laser-doped lines which may be formed during the laser doping process, can provide further alterative paths for the penetration of Cu into the Si substrate.

While using EDX spectrometry in TEM was shown to be an effective method to detect the diffused Cu in Si, this method suffers from being tedious, costly and having a

73 relatively low detection limit. Furthermore, using the TEM system for quantitative analysis is problematic and unreliable. These disadvantages necessitate complimentary characterisation techniques to be used in parallel with TEM.

Depth profiling techniques such as secondary ion mass-spectrometry (SIMS) and AES have been widely used to measure the diffusion profile. However, these tools are also costly and time consuming and typically not very effective on textured surfaces. Electrical measurement tools on the other hand, have proved to be more sensitive than depth profiling techniques in characterisation of Cu diffusion barriers (Istratov et al. 2000b, Istratov and Weber 2002). Although still qualitative, electrical measurement techniques have the advantage of being able to detect lower concentrations of Cu contamination in solar cells, than is possible with EDX/TEM. Furthermore, for this study it is critical to understand the effect of diffused Cu in the Si on the electrical properties of solar cell devices.

74 Chapter 4 Impact of Copper Contamination on the Electrical Properties of Laser-Doped Selective- Emitter Solar Cells

4.1 Introduction

In Chapter 3, EDX was used in conjunction with TEM as a direct method for detecting Cu which had penetrated through electroless Ni and light-induced plated Ni diffusion barrier layers in LDSE solar cells. The results of chemical identification of Cu in Si obtained from the EDX/TEM analysis suggested that Ni barrier layers fail to block Cu from diffusing into Si. Most of the Cu precipitates were observed in the highly-doped selective-emitter region under the laser-doped fingers (the n++ area). If Cu diffuses out of these regions and forms precipitates in the adjacent lightly-doped emitter layer (the n+ area) or along the p-n junction area, the cell performance would be expected to degrade dramatically. Given that these precipitates are most probably in the form of a

Cu-rich silicide (Cu3Si), they form a defect energy level close to the middle of the Si bandgap. This will reduce the lifetime of minority carriers by introducing effective recombination sites (Istratov et al. 2000a, Istratov and Weber 2002) and severely affect the electrical properties of solar cells. Furthermore, precipitates that span the p-n junction will result in localised shunting of the junction.

The proposed replacement of Ag screen-printed metallisation with Cu plating for front side metallisation of Si solar cells makes it technologically a critical task to evaluate the long term effects of Cu on the electrical properties of solar cells. Indeed, electrical characterisation techniques are reported to be more sensitive than depth profiling techniques or EDX/TEM in detecting low concentrations of Cu in Si (Istratov et al. 2000b, Istratov and Weber 2002). 75 This chapter reports an investigation into the effects of Cu contamination on the electrical performance of LDSE solar cells. For this purpose, a simple, fast and more effective characterisation method is required. Measurement of current-voltage (I-V) curves and high resolution PL imaging are the main device characterisation techniques employed in this thesis. The electrical analysis methods employed in this work are qualitative, whereas the PL imaging method is used to quantitatively determine the recombination current of damaged regions in LDSE solar cells. Both analysis methods have the benefit of being non-destructive to the sample while being fast and providing high throughput.

To prompt the Cu diffusion from the front side metal contacts into the Si bulk of LDSE solar cells, samples were exposed to heat treatments of various temperatures and durations. The efficiency (η) and the local ideality factor (m) of the cells obtained from illuminated and dark I-V curves, respectively, along with the PL images were measured for each sample before and after heat treatment.

Background information on the I-V measurements and analysis is presented in Section 4.2.1. Then, Section 4.2.2 describes the PL method used to determine quantitative values for the recombination current. Section 4.3 contains a description of the experimental process for this study. The results from the investigation into the electrical properties of the cells are presented and discussed in Section 4.4 and Section 4.5, respectively.

4.2 Theory of Analysis Methods

4.2.1 Characterisation of Current-Voltage Curves

Solar cells are typically characterised and compared by measuring their efficiency. The different parameters that affect the efficiency of a solar cell are commonly revealed by the characteristic I-V curves of the cell. The I-V curve is the relationship between the output current of the solar cell I and the operating voltage of the cell V. The I-V curve can be measured with the cell under illumination (illuminated I-V curve) or in the dark (dark I-V curve). A solar

76 cell‟s I-V curve is generated by varying the cell voltage and plotting the measured corresponding cell current.

The illuminated I-V curve reveals the primary parameters affecting the efficiency of the solar cell. The open circuit voltage Voc, the short circuit current Isc and the fill factor FF determine the efficiency, , by (Wenham et al. 2006a):

퐼푠푐 . 푉표푐 . 퐹퐹 휂 = (2.6) 푃𝑖푛

Where Pin is the power of the light incident on the cell. The open circuit voltage Voc is the maximum voltage of the cell at zero current, the short circuit current Isc is the maximum current of the cell at zero voltage and the FF is a measure of the “squareness” of the I-V curve and determines the maximum power from the solar cell. The FF is defined as:

퐼푚푝푝 . 푉푚푝푝 퐹퐹 = (2.7) 퐼푠푐 . 푉표푐

Impp and Vmpp are the current and voltage at maximum power point (MPP), where the power output of the solar cell, i.e., the product of I and V is at its maximum value. Both FF and efficiency are typically used to access the performance of solar cells. Figure 4.3a displays a typical illuminated I-V curve of a LDSE solar cell, showing the open circuit voltage, the short circuit current and the current and voltage at MPP.

Solar cells are frequently simulated using a “two-diode model” with the equivalent electrical circuit as shown in Figure 4.1 (McIntosh 2001, Kunz 2009). This model consists of a light-generated current source IL in parallel with two diodes which represent the p-n junction, plus the parasitic series and shunt resistances, Rs and Rsh respectively. The currents flowing across the two diodes are known as diodes‟ dark current and are given by Shockley‟s ideal diode equation (Shockley 1949):

푉푑 퐼1 = 퐼01 푒푥푝 − 1 (2.8) 푛1푉푇

77 and

푉푑 퐼2 = 퐼02 푒푥푝 − 1 (2.9) 푛2푉푇

where VT is the “thermal voltage”: 푘푇 푉푇 = 푞 = 25.8 푚푉 푎푡 300 퐾 (2.10) q is the charge of an electron, k is the Boltzmann‟s constant, T is the absolute temperature, Vd is the voltage across the p-n junction, I01 and I02 are the dark saturation currents of the two diodes, respectively and n1 and n2 are the ideality factors of the two diodes (n1=1 for the first diode and n2=2 for the second diode).

Rs I1 I2 IRsh

IL Rsh V

Figure 4.1: The two-diode equivalent circuit model of a solar cell under illumination.

I01 is related to the recombination in the bulk and at the exposed surfaces of the solar cell, whereas I02 is caused by recombination in the depletion-region of the p-n junction and at the edges of the solar cell (Green 1998, McIntosh 2001, Kunz 2009). An increase in the recombination rates and therefore the saturation currents of the solar cell leads to a reduction in the Voc of the cell. Recombination in the depletion region and therefore the larger I02 can also reduce the FF (Green 1998, McIntosh 2001).

The shunt resistance Rsh is characteristic of a low resistance path parallel to the p-n junction which causes current leakage across the junction. The main contributors to Rsh are non-ideal junction properties, crystal defects, impurities near the junction and manufacturing defects. The series resistance Rs is the electrical resistance of a solar cell and is due to the bulk resistance of the Si wafer, the resistance associated with the lateral current flow in the emitter, the resistance of the metal contacts and the cell

78 interconnection metal and the contact resistance between the metallic contacts and the

Si wafer of solar cells (Green 1998, Wenham et al. 2006a). Both Rs and Rsh reduce the FF of the solar cell (Green 1998, McIntosh 2001).

As can be seen from Figure 4.1, the output current of the solar cell is:

퐼 = 퐼 − 퐼 − 퐼 − 퐼 퐿 1 2 푅푠ℎ (2.11)

The illuminated I-V characteristic of the two-diode model solar cell is obtained by measuring the I-V curve of the cell under constant illumination. Accordingly, Equation

4.6 which includes the voltage drop across Rs and the current flow through Rsh can be defined as:

푉푑 푉푑 푉푑 퐼 = 퐼퐿 − 퐼01 푒푥푝 − 1 − 퐼02 푒푥푝 − 1 − (2.12) 푛1푉푇 푛2푉푇 푅푠ℎ and

푉푑 = 푉 + 퐼 × 푅푠 (2.13)

In order to measure the dark I-V curve, a voltage is applied across the solar cell in the dark. In this case IL=0. While the main parameters affecting the performance of a solar cell can be determined from the light and dark I-V curves, the local ideality factor curve (m-V curve) reveals information that may be subtle in the corresponding I-V curves. The local ideality factor (m) can be derived from the illuminated and dark I-V curves and relates to the local slope ln(I)-V curve of a solar cell defined as (McIntosh 2001):

1 푑푉 퐼 푑푉 푚 = . = . (2.14) 푉푇 푑 푙푛퐼 푉푇 푑퐼

Although the two-diode model is often used to simulate the solar cell characteristics, the LDSE solar cells exhibit a different behaviour which does not conform to the two-diode equivalent electrical circuit. McIntosh suggested a modified circuit with an additional circuit element which is comprised of a diode and a resistor as shown in Figure 4.2 (McIntosh 2001). This additional part represents a region with high recombination rate in a solar cell and is known as a resistance-limited enhanced recombination region.

79 I

RH Rs IH I1 I2 IRsh

IL Rsh V

Figure 4.2: The equivalent circuit representation suggested by McIntosh for a solar cell that contains a region of resistance-limited enhanced recombination, under illumination (McIntosh 2001).

The high recombination current flowing through the diode in this region is given by:

푉퐻 퐼퐻 = 퐼0퐻 푒푥푝 − 1 (2.15) 푛퐻푉푇

noting the diode is separated from the main part of the solar cell by a resistor RH (McIntosh 2001). The source of the high recombination rate in LDSE solar cells is associated with the formation of a Schottky diode contact which is due to the penetration of metal into the lightly-doped area (Tjahjono et al. 2007a).

Therefore, the I-V characteristic expressed in Equation 4.7 is modified to:

푉푑 푉푑 푉퐻 푉푑 퐼 = 퐼퐿 − 퐼01 푒푥푝 − 1 − 퐼02 푒푥푝 − 1 − 퐼0퐻 푒푥푝 − 1 − 푛1푉푇 푛2푉푇 푛퐻푉푇 푅푠ℎ (2.16) where

푉퐻 = 푉푑 − 퐼퐻 × 푅퐻 (2.17)

The m-V curve of such solar cells contains a “hump” with an ideality factor greater than unity between low to moderate voltages. For clarification, the I-V characteristics of a typical LDSE solar cell which exhibits the m-V hump are depicted in Figure 4.3b.

80 a) b)

0.30 1.E+00 1.E-01 0.25 1.E-02 Isc V , I mpp mpp 1.E-03 0.20 1.E-04

0.15 1.E-05

Current (A) Current 1.E-06 Current (A) Current 0.10 1.E-07

Voc 1.E-08 0.05 1.E-09 5.0 0.00 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 4.0 Voltage (V)

3.0

2.0

1.0 Local Ideality Factor Ideality Local

0.0 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 Voltage (V)

Figure 4.3: a) Illuminated I-V curve of an example of a LDSE solar cell, showing the open circuit voltage Voc, short circuit current Isc and current and voltage at maximum power point (Impp and Vmpp); b) corresponding dark I-V and m-V curves exhibiting a hump in the local ideality factor curve at low to moderate voltages.

The magnitude of I0H affects the peak-width of the hump, i.e., as I0H increases the increasing slope begins at lower voltages. For I0H values that are equal or smaller than

I01, no hump is evident in the m-V curve which means that there is no significant region of high recombination. However, if I0H is greater than I01, the m-V curve exhibits a hump which leads to a reduced FF. On the other hand, the magnitude of RH affects the voltage at which the hump is evident. As RH decreases, the hump is shifted to a larger voltage value and thus can reduce the efficiency of the solar cell. In summary, the detrimental effect of the hump can be avoided if RH is large enough so that the m-V hump occurs at a significantly lower voltage than the MPP voltage or if I0H is smaller than I01.

Since the hump in the local ideality factor curve is commonly observed in LDSE solar cells, for the I-V analysis of the experimental work reported in this thesis, the “hump” model was used to model the LDSE cells. As mentioned earlier, the diffusion of Cu from the front side metal contacts into the lightly-doped emitter in LDSE cells is

81 detrimental to minority carrier lifetime because Cu precipitates in Si act as effective recombination centres (Weber 1983, Baumann et al. 1997, Istratov et al. 1998, Istratov and Weber 1998, Seibt et al. 1998, Istratov et al. 2000b, Buonassisi et al. 2005). Furthermore, localised shunting and/or Schottky diode contacts may form if Cu penetrates through the lightly-doped emitter and contacts the underlying p-type Si in localised regions. These Schottky contacts are the most probable causes of the m-V hump in LDSE cells. Therefore, examination of the I-V and m-V curves can reveal valuable information about the sources of high-recombination regions and Schottky contacts. The I-V analysis presented in this chapter examines the effect of Cu diffusion on the hump characteristics in the local ideality factor curve. Photoluminescence imaging, on the other hand, was used to investigate shunts in LDSE solar cells which may be created or worsened by Cu precipitation in Si.

4.2.2 Photoluminescence Imaging

Photoluminescence imaging is a fast and powerful spatially-resolved characterisation technique for Si wafers and Si solar cells. The data acquisition time for the PL technique is very short, on the order of seconds, or less. Also, being a contactless measurement technique, it allows the non-destructive analysis of solar cells at any individual stage of the fabrication process (Trupke et al. 2006a, Trupke et al. 2006c).

Luminescence is the emission of light (photons) from a semiconductor under external excitation. The different types of luminescence are determined according to the method of excitation. When photon emission is produced as a result of exposing the semiconductor to a light source, the process is called PL. Electroluminescence (EL) refers to the process where photon emission is caused by injection of an electric current.

Once excited, the electrons in the semiconductor have enough energy to jump from a lower energy valence band (VB) to a higher energy conduction band (CB), leaving behind a hole. The electrons in the CB eventually lose their energy and recombine with holes in the VB. This process is an energy loss mechanism and is known as recombination. Recombination in a semiconductor can occur radiatively and non- radiatively. The main two types of non-radiative recombination are Auger and Shockley-Read-Hall (SRH) recombination.

82 In Auger recombination, an electron in the CB recombines with a hole in the VB, but instead of emitting the energy as heat or photons, the excess energy is given to a second electron in the CB. This second electron eventually releases the energy in the form of phonons (Green 1998).

Shockley-Read-Hall recombination is a two-step process which occurs due to the presence of defect levels. Defects in a semiconductor can be introduced by impurities (e.g., metals) or crystallographic defects (e.g., grain boundaries, surfaces, point defects). These defects introduce discrete energy levels within the otherwise forbidden bandgap. An electron in the CB is first trapped by the energy state in the forbidden bandgap, before recombining with a hole in the VB. The energy is usually given up as phonons (Green 1998).

PL imaging is based on the measurement of radiative recombination. In radiative recombination, an electron in the CB recombines directly with a hole in the VB losing all or part of its energy as a photon. The energy of the emitted photon is similar to the bandgap energy. Photons which are emitted from the sample are detected and measured as a PL signal. The PL signal is determined by the radiative recombination rate which is proportional to the product of minority and majority carrier concentrations. Therefore, the PL signal is given by (Trupke and Bardos 2005a, Trupke and Bardos 2005b):

푃퐿 = 퐴𝑖. 퐵. 푛푒 . 푛ℎ = 퐶 휏 . 푛푒 . 푛ℎ (2.18)

where B is the radiative recombination coefficient, Ai accounts for the specific optical properties of the sample, ne and nh are the electron and hole concentrations, respectively. C(τ) is thus a proportionality constant that accounts for the sample lifetime and optical properties (Green 1998, Trupke and Bardos 2005a).

At any point within the sample, the carrier concentrations are related to the splitting of the quasi-Fermi energies (Würfel and Ruppel 1981, Green 1998):

2 푛푒 . 푛ℎ = 푛𝑖 . 푒푥푝(∆퐸퐹 푘푇) (2.19)

83 where ni is the intrinsic carrier concentration and ΔEF is the splitting of the quasi-Fermi levels at the junction (Würfel and Ruppel 1981, Green 1998). Under steady state excitation, ΔEF can be related to the junction voltage Vd by (Würfel and Würfel 2009):

∆퐸퐹 = 푞. 푉푑 (2.20)

Considering the relationship given in Equation 4.15, then combining Equation 4.13 and Equation 4.14, the PL signal is given by (Würfel and Ruppel 1981, Würfel 1982):

2 푃퐿 = 퐶 휏 . 푛𝑖 . 푒푥푝(푉푑 푉푇) (2.21) or:

푃퐿 = 퐶. 푒푥푝(푉푑 푉푇) (2.22)

2 where 퐶 = 퐶 휏 . 푛𝑖 . In an open circuit PL image, PL is the average PL signal from the sample. In a Si wafer, Vd is the implied voltage and in a finished solar cell, Vd is equal to the measurable junction voltage (Augarten et al. 2009) and VT is the thermal voltage (25.8 mV at 300 K).

In the PL imaging system used in this thesis, an 815 nm diode laser was used to uniformly illuminate the entire sample area. The PL signal emitted by the sample was captured by a Si charge-coupled device (CCD) camera cooled to -30 ° C. The spatial resolution is typically 165 µm per pixel. However, for the experimental work in this thesis, a zoom lens was used to give a higher image resolution of 30 µm per pixel.

PL imaging methods have been developed to quantitatively determine several cell parameters. Of particular relevance to the experimental work of this thesis is shunting because it could be caused by the precipitation of Cu close to the junction in the LDSE cells. Shunting can result in reduced Voc and FF values (Green 1998). PL imaging has been previously demonstrated as an efficient tool to qualitatively detect shunts (Trupke et al. 2006b). More recently, a method of quantitatively determining shunt resistance values in Si solar cells using an open circuit PL image has been developed by Augarten (Augarten et al. submitted January 2011). An approach based on this method was used in the experimental work of this thesis to quantitatively analyse process-induced changes in the current flowing through shunt paths. A brief description of this method is presented below. 84 4.2.2.1 Photoluminescence Imaging for Quantitative Shunt Calculations

As described in Section 4.2.1‎ , shunts are generally modelled in solar cells as low resistance paths parallel to the junction. Shunts in Si solar cells are often unintentionally created during processing. For example, during sintering of plated metal contacts, metal may penetrate into the Si and form a low-resistance pathway between the n-type and the p-type regions of the junction. The shunt resistance Rsh is given as:

푉푅푠ℎ 푅푠ℎ = (2.23) 퐼푅푠ℎ

where VRsh is the local voltage at the shunt and IRsh is the current flowing through the shunt path. Using an open circuit PL image, both VRsh and IRsh can be calculated, as explained below (Augarten et al. 2009).

 Calculation of the Local Voltage at the Shunt

In order to calculate VRsh from an open circuit PL image, the calibration constant C 푉 should be first obtained from Equation 4.17, (푃퐿 = 퐶. 푒푥푝 푑 ), where PL is the 푉푇 average PL intensity taken from the entire cell area and Vd is the measured Voc of the sample. Once C is known (C is assumed to be constant (Augarten et al. 2009)), VRsh can be calculated using the same equation. To do so, Vd is replaced by VRsh and the lowest PL signal at the site of the shunt is used for the PL term. The PL signal is proportional to the effective minority carrier lifetime. Consequently, the darkest regions in the PL open circuit image (i.e., the regions with lowest PL signal), correspond to regions of low lifetime which can be caused by recombination centres or shunts in solar cells (Augarten et al. 2009).

 Calculation of the Shunt Current

A local shunt in an illuminated solar cell acts as a local current sink, i.e., a proportion of the current from the surrounding non-shunted cell area flows into the shunt. The shunt current IRsh is the total current that is extracted from the surrounding non-shunted area. In a PL open circuit image, the non-shunted area around a local shunt appears as a

85 circular blurred region of reduced luminescence which is due to the effect of the shunt on the surrounding region (Augarten et al. 2009).

To calculate IRsh, a “one-diode” model is used to describe the equivalent electrical circuit of a shunted region of a solar cell and the area affected by the shunt. Figure 4.4 depicts a shunt region and an associated equivalent circuit. For purposes of illustration the shunt is shown as a dark area at the centre, surrounded by a number of concentric rings, each with an area ai. These rings represent unshunted regions, however current flowing from these regions is affected by the shunt located at the centre. The equivalent electrical circuit model is shown in Figure 4.4b as a number of series connected one- diode models. The diode dark current density in each area is given by Shockley‟s ideal diode equation (Shockley 1949) as:

푉푑,𝑖 퐽푑,𝑖 = 퐽01,𝑖. 푒푥푝 (2.24) 푛1푉푇

where Vd,i is the diode voltage, Jo1,i is the dark saturation current density of each area and n1 is the ideality factor. For this method a value of n1=1 is adopted, which is characteristic of an ideal diode in which recombination occurs predominantly in the bulk and at surfaces (Augarten et al. 2009).

86 a)

b) Node 1 Node i

IL,1 IRs,1 IL,i IRs,i

Id,1 IRsh Id,i

Figure 4.4: a) The circular representation of a shunted region in a solar cell with the shunt in the centre surrounded by the unshunted area with reduced luminescence; b) the equivalent electrical circuit using the one-diode model (Augarten et al. submitted January 2011).

87 The light generated current density JL is assumed constant across the cell. The generated current in each area ai is:

퐼퐿,𝑖 = 퐽퐿. 푎𝑖 (2.25)

As explained above, a proportion of the current in each area is extracted by the shunt. As can be seen in Figure 4.4b, the extracted current density which flows into the shunt from each area ai is given by (Augarten et al. 2009):

퐽푒푥푡푟 ,𝑖 = 퐽퐿 − 퐽푑,𝑖 (2.26)

The current flowing into the shunt IRsh is the total extracted current as:

퐼푅푠ℎ = 퐼푒푥푡푟 ,𝑖 𝑖 or: (2.27)

퐼푅푠ℎ = 퐽푒푥푡푟 ,𝑖. 푎𝑖 𝑖

On the other hand, it is indicated from Equation 4.17 that the PL signal is related to the diode voltage. Therefore, combining Equation 4.17 with the Shockley‟s ideal diode law given in Equation 4.19, and assuming J01 is constant throughout the area of the cell, the local PL signal from any area ai can be written as (Augarten et al. 2009): 퐶 퐶 푃퐿𝑖 = . 퐽푑,𝑖 = . 퐽퐿 − 퐽푒푥푡푟 ,𝑖 (2.28) 퐽01 퐽01

It can be seen from Equation 4.23 that an increase in the amount of current extracted by the shunt at each area ai (Jextr,i) leads to a reduction in the PL signal from that area.

Therefore, it can be deduced that the shunt current IRsh is proportional to the total reduction in the PL signal due to the shunt (Augarten et al. 2009). Figure 4.5a presents an example of a PL image of a solar cell with a shunt. The PL profile across the shunt is shown in Figure 4.5b. The edges of the shunted region in the PL profile represent the area which is not affected by the shunt. The shaded region is the difference between the PL signals of the shunted and the non-shunted regions and is proportional to the shunt current IRsh (Augarten et al. submitted January 2011).

a) b)

Iextr,i PLnon-shunted

Reduction in PL signal due to the shunt PL Signal PL PLi Shunt Distance from Shunt

Figure 4.5: a) A PL image of a solar cell with a shunt; b) profile of the PL signal across the shunt is shown as the black curve. The red dotted curve is the theoretical non- shunted PL signal. The shunt current IRsh is proportional to the difference between the PL signals of the shunted and the non-shunted regions (the shaded area) (Augarten et al. submitted January 2011).

The PL signal in the non-shunted area can be obtained from Equation 4.23, by setting

Jextr,i=0 as no current is extracted by the shunt. The PLnon-shunted is therefore given by (Augarten et al. 2009): 퐶 푃퐿푛표푛 −푠ℎ푢푛푡푒푑 = . 퐽퐿 (2.29) 퐽01

where 퐽퐿 = 퐼퐿 퐴푐푒푙푙 . IL is the total light generated current under the illumination intensity during the measurement, and Acell is the total area of the cell. The illumination wavelength used in this thesis is 815 nm, so it is assumed that light not reflected by the sample is absorbed. Therefore, JL is given by:

퐽퐿 = 푞. 휙. 1 − 푅 (2.30) where q is the electron charge, Ф is the incident photon flux (photons/s/cm2) and R is the reflectance of the sample at 815 nm.

Equation 4.24 can be written as:

89 퐶 푃퐿푛표푛 −푠ℎ푢푛푡푒푑 = (2.31) 퐽01 퐽퐿

According to Equation 4.26, measuring the PL intensity in an area of the cell not 퐶 affected by the shunt (PLnon-shunted) allows determination of the ratio: . Re- 퐽01 arranging Equation 4.23 leads to:

푃퐿𝑖 퐽푒푥푡푟 ,𝑖 = 퐽퐿 − 퐶 (2.32) 퐽01

Jextr,i can be determined by combining Equation 4.26 and Equation 4.27 (Augarten et al. 2009):

푃퐿𝑖 퐽푒푥푡푟 ,𝑖 = 퐽퐿. 1 − (2.33) 푃퐿푛표푛 −푠ℎ푢푛푡푒푑

It can be seen from Equation 4.28 that by measuring the local PL signal (PLi) at every pixel in the PL image deemed to be affected by the shunt, Jextr,i can be calculated for each pixel. The extracted current can subsequently be calculated upon multiplication by ai, which can be taken as the area of a single pixel of the PL image:

푃퐿𝑖 퐼푒푥푡푟 ,𝑖 = 퐽퐿. 1 − . 푎𝑖 (2.34) 푃퐿푛표푛 −푠ℎ푢푛푡푒푑

Finally, the total current flowing into the local shunt is the sum of the local current extracted from all pixels in the area affected by the shunt:

푃퐿𝑖 퐼푅푠ℎ = 퐽퐿 . 1 − . 푎𝑖 푃퐿푛표푛 −푠ℎ푢푛푡푒푑 (2.35) 𝑖

In summary, IRsh can be calculated using the value of JL determined from

Equation 4.25, the PL signal in the area not affected by the shunt (PLnon-shunted), the local PL signal in the area affected by the shunt (PLi) and the area ai all of which can be obtained from an open circuit PL image using an image processing and analysis program such as ImageJ (ImageJ 2011).

The shunt resistance can therefore be obtained from Equation 4.18, with VRsh and IRsh calculated from an open circuit PL image.

90 The procedure described above for calculation of the shunt resistance is applicable ideally to shunt regions that do not overlap with metal fingers in a finished solar cell. Metallised regions are evident in PL images because they shade the underlying Si, and are typically sites of high carrier recombination. As a result, metal regions appear as dark areas in PL images. Furthermore, the low resistance of the metal grid alters the direction of current flow from nearby regions (Kasemann et al. 2007). Therefore, if a shunt is located close to the front metallised region, the current extraction through the metal grid reduces the effect of the shunt on the local PL signal. Consequently, it is difficult to identify the boundaries of shunt regions that are located close to metal lines.

In the current work, shunt regions potentially caused by Cu precipitates occur in the vicinity of metal fingers, so a slightly different approach was used to analyse the PL images. To quantitatively analyse the changes in the current flowing through local shunts, the following steps were performed:

 An open circuit PL image of the finished cell was captured, under an appropriate laser voltage, exposure time and area setting.  The image was spatially calibrated using the ImageJ program. It is important to define the spatial scale of the images so that the measurements are according to the real cell dimensions. Since a zoom lens was used in the PL imaging system for the current experiments, the corresponding spatial scale of approximately 33 pixels per mm was used for all experimental work in this thesis.  The gray scale level (luminescence information) was calibrated to make sure all images are compared using the same scale.  A section containing a low PL signal conceivably due to a local shunt was identified. Then, an area around the shunt was selected using the rectangular area selection tool. Note the horizontal distance of the selected analysis area for all quantitative shunt calculations in this chapter is equal to the metal finger spacing which is 1 mm.  An example of the open circuit PL image of a finished solar cell and the selected area is shown in Figure 4.6.  A two-dimensional graph of the PL intensities of pixels along the selected area was obtained. This plot profile was generated using the “Plot” command in the

91 “Analyse” menu in ImageJ program. The X-axis of the plot represents the horizontal distance through the selection and the Y-axis displays the gray value, i.e., the average intensity over the vertical column of pixels. Figure 4.6 shows the corresponding plot profile of the selected rectangular area. As can be seen, point “a” and point “b” display the two maximum values of the PL intensity obtained from the plot profile of the selected area. The PL signal in the area not

affected by the shunt (PLnon-shunted) is determined by averaging the PL intensity values of these two points:

푃퐿푎 + 푃퐿푏 푃퐿 = 푛표푛 −푠ℎ푢푛푡푒푑 2 (2.36)

PLa PLb

1 mm

Figure 4.6: An example of the open circuit PL image of a finished solar cell and the selected rectangular area for analysis. The corresponding plot profile of the selected area is also shown, exhibiting point “a” and point “b” as the two maximum values of the PL intensity.

92  The “mean” gray scale value of the selected area was determined, i.e., the average PL intensity of all pixels in the selected area. Using the “Measure” command in the “Analyse” menu in ImageJ program, the average PL intensity of the whole selected area was calculated and recorded as:

𝑖 푃퐿𝑖 푚푒푎푛 = 푁 (2.37)

where N is the number of pixels in the selected area. Finally, IRsh can be calculated by re-writing Equation 4.30 as: 1 푃퐿 퐼 = 퐽 . 푎 . 푁. 1 − . 𝑖 𝑖 푅푠ℎ 퐿 𝑖 (2.38) 푃퐿푛표푛 −푠ℎ푢푛푡푒푑 푁

푃퐿𝑖𝑖 then substituting PLnon-shunted and 푁 determined from Equation 4.31

and Equation 4.32, respectively.  An open circuit PL image from the cell was captured after heat treatment under the same laser voltage, exposure time and area setting.

 The above steps were repeated to calculate IRsh in the same selected area of the cell after heat treatment.

 The difference in the calculated IRsh values of the selected area before and after heat treatment was considered indicative of the increase in the current flowing through the local shunt.

Note that in all the PL analysis reported in this chapter, the selected analysis area contained a part of the metal finger. As shown in the plot profile in Figure 4.6, the PL intensity over the metal finger is minimum since the front metal grid blocks the PL signal from underneath it. Also, the metal presents a low-resistance path to the surrounding area and therefore, reduces the effect of the local shunt on the PL signal. This causes errors in calculating absolute values of the shunt resistance close to the metallised region. Nevertheless, the influence of the shunt extends along the area next to the metal finger and consequently reduces the PL intensity from the affected area. As a result, by comparing the reduction in the PL intensity from the same analysis area before and after heat treatment, and calculating the increase in the current flowing into the shunt after heat treatment, a relative estimation of the increase in the severity of the

93 shunt can be deduced. This effect is presumably caused by precipitation of Cu in Si close to the junction in the LDSE cells.

4.3 Experimental

Two groups of LDSE solar cells were fabricated as described in Section 3.2.1. The Ni barrier layer in the first group of cells was deposited by electroless plating whereas for the second group the Ni barrier layer consisted of LIP Ni. Section 3.2.1.1 and Section 3.2.1.2 provide a detailed description of the metallisation process used for both groups.

The finished cells from each group were then paired in sub-groups for a subsequent heat treatment. Each sub-group consisted of one cell with a Ni layer formed by electroless plating, and a second cell with a LIP Ni layer. Table 4.1 summarises the heat treatment parameters and cooling method used for sintering the sub-groups. All cells were heat- treated in a muffle oven. A fast cooling procedure of quenching in ethylene glycol was used for the groups QA to QD. To investigate the effect of cooling rate on the electrical properties of the cells, two sub-groups of SA and SB were slowly cooled under ambient laboratory conditions.

To be consistent with the experimental procedure employed for chemical microanalysis of Cu diffusion in Chapter 3, two different temperatures of 400 °C and 200 °C were used for heat treating the cells prior to device characterisation. The cells in the first sub- group, QA, were used for a test experiment to examine the implications of a 5 hour of thermal stress at 400 °C on the electrical properties of the cells. These cells were subsequently quenched in ethylene glycol. It was found that a thermal treatment under these conditions was highly detrimental to cell performance and in some cases, made it impractical to analyse the electrical properties of the cells after heat treatment. Therefore, the temperature was scaled down to 200 °C to heat treat the majority of the cells. In addition, the sintering temperature of 200 °C was appropriate for the electrical analysis of the solar cells since it is closer to the typical thermal stress that the solar cells would endure under outdoor conditions.

The cells in the sub-groups of QB to QD underwent heat treatments at 200 °C for various durations of 1.30, 2.0 and 2.30 hours, respectively, followed by quenching in 94 ethylene glycol. The cells from the remaining two sub-groups SA and SB were exposed to a thermal stress at 200 °C for 2 and 2.30 hours, respectively. These cells were cooled at a slow rate under ambient laboratory conditions.

All cells were characterised before and after heat treatment by measuring and comparing their dark and illuminated I-V curves using an I-V tester known as “Darkstar”. Additionally, a luminescence imaging system from BTimaging Inc.4 with a zoom lens was used to capture open circuit PL images for all cells before and after heat treatment. Results from the I-V characterisation and PL imaging analysis are presented in Section 4.4.1 and Section 4.4.2, respectively.

Table 4.1: Summary of heat treatment parameters and cooling method.

Heat Treatment Group Cooling Method Temperature (°C) Time (hr:mm)

QA 400 5:00 Quenched QB 200 1:30 Quenched QC 200 2:00 Quenched QD 200 2:30 Quenched SA 200 2:00 Slow Cooled SB 200 2:30 Slow Cooled

4.4 Results

4.4.1 Efficiency and I-V Characterisation

The illuminated and dark I-V and m-V curves presented in Figure 4.7 and Figure 4.8 were measured for the cells in group QA, with an electroless Ni barrier layer and a LIP Ni barrier layer, respectively. Each figure contains the I-V characteristics before and after a heat treatment consisting of a 5 hour of thermal stress at 400 °C and subsequent quenching in ethylene glycol. It is clear from the I-V curves that the Voc and the efficiency of both cells decreased sharply as a result of the heat treatment.

4http://www.btimaging.com/?page_id=81 95 For the cell with electroless Ni as the diffusion barrier layer, the m-V hump exhibited by the cell is exacerbated after the heat treatment. Figure 4.7b displays the increase in the height of the hump from 4.3 to 9.3. Furthermore the hump peak position shifted to a larger voltage and the full width at half maximum (FWHM) of the hump increased by

25%. As a result, the hump extended beyond the MPP voltage and the Voc. As explained in ‎4.2.1, the m-V hump can cause a reduction in the FF and the Voc if it extends beyond the MPP voltage. In this case the FF reduced from 0.743 to 0.492, coinciding with a reduction in the absolute cell efficiency from 17.45% to 5.61%.

The other cell from group QA which used LIP Ni as the diffusion barrier layer, exhibited a similar behaviour to the cell with electroless Ni as the barrier layer. It can be seen from Figure 4.8 that the 5 hour heat treatment at 400 °C and a subsequent quenching cause the FF and the Voc of the cell to be remarkably reduced. The absolute cell efficiency decreased from 15.92% to 4.68%. However after heat treatment the m-V hump could no longer be identified. Instead the value of m increased monotonically with increasing cell voltage (not shown). A change in the m-V characteristic of this nature was typically observed to coincide with a severe reduction in cell efficiency.

As observed from the I-V analysis of the cells in group QA, both cells with electroless Ni and LIP Ni barrier layers suffered a significant performance loss after the 5 hour heat treatment at 400 °C followed by quenching. Due to such a severe degradation of cell performance it is difficult to establish whether the variations in the I-V characteristics between the two barrier layers were significant. However, for a less severe thermal stress and subsequent quenching, the I-V analysis revealed that the type of Ni barrier layer has a significant impact on the final electrical performance of the cell.

96 a) b)

Before HT After HT Before HT After HT

0.35 1.E+01 0.30 1.E+00 1.E-01 0.25 1.E-02 1.E-03 0.20 1.E-04 1.E-05 0.15 Current (A) Current 1.E-06 Current (A) Current 1.E-07 0.10 1.E-08 1.E-09 0.05 1.E-10 10 0.00 9 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 8 Voltage (V) 7 6 Before heat treatment: After heat treatment: 5 Voc= 635 mV Voc= 322 mV 4 2 2 Jsc= 37 mA/cm Jsc= 35.3 mA/cm 3

FF= 0.743 FF= 0.492 ideality Factor Local 2 η= 17.45 η= 5.61 1 0 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 Voltage (V) Figure 4.7: The I-V characteristics of the cell from group QA with electroless Ni as the barrier layer, before and after 5 hours of heat treatment at 400 °C, followed by quenching in ethylene glycol, a) the illuminated I-V curve, b) the corresponding dark I-V and m-V curves.

a) b)

Before HT After HT Before HT After HT

0.35 1.E+01 0.30 1.E+00 1.E-01 0.25 1.E-02 1.E-03 0.20 1.E-04 1.E-05 0.15 Current (A) Current 1.E-06 1.E-07 0.10 Current (A) Current 1.E-08 1.E-09 0.05 1.E-10

0.00 5 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 4 Voltage (V)

3 Before heat treatment: After heat treatment: Voc= 634 mV Voc= 303 mV 2 2 2 Jsc= 39 mA/cm Jsc= 35 mA/cm 1 FF= 0.651 FF= 0.441 ideality Factor Local η= 15.92 η= 4.68 0 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 Voltage (V)

Figure 4.8: The I-V characteristics of the cell from group QA with LIP Ni as the diffusion barrier layer, before and after 5 hours of heat treatment at 400 °C, followed by quenching in ethylene glycol, a) the illuminated I-V curve, b) the corresponding dark I-V and m-V curves.

97 Figure 4.9 illustrates the observed changes in cell efficiency following the various heat treatments, for both electroless Ni and LIP Ni diffusion barriers. The efficiency of the cells featuring electroless Ni as the diffusion barrier layer decreased markedly as a result of heat treatment for each temperature-duration combination. In comparison, the cells with LIP Ni as the barrier layer show a far smaller drop in efficiency when heated at 200 °C (groups QB to QD).

18.00 17.00 LIP Ni/LIP Cu 16.00 Before HT 15.00 14.00 13.00 12.00 LIP Ni/LIP Cu 11.00 After HT 10.00 9.00 8.00 Electroless Ni/LIP Cu 7.00 Before HT

Efficiency (%) Efficiency 6.00 5.00 4.00 3.00 Electroless Ni/LIP Cu 2.00 After HT 1.00 0.00 Group QB Group QC Group QD Group QA

1:30 hrs at 200 C 2:00 hrs at 200 C 2:30 hrs at 200 C 5:00 hrs at 400 C Quenched Quenched Quenched Quenched

Figure 4.9: The effect of various heat treatments followed by a quench in ethylene glycol on the absolute efficiency of LDSE solar cells. The cells with electroless Ni and LIP Ni as the diffusion barrier layer are marked by green and orange, respectively.

This trend can also be observed by visually inspecting the m-V curves of the cells. Figure 4.10 and Figure 4.12 provide examples of the I-V and m-V curves of representative cells from group QC, before and after the 2 hour of thermal stress at 200 °C, followed by a quench in ethylene glycol, with electroless Ni and LIP Ni barrier layers, respectively. It can be seen in Figure 4.10b that the m-V hump of the cell with electroless Ni differed significantly before and after thermal treatment. In contrast, the m-V hump of the cell with LIP Ni was less affected upon heat treatment and subsequent quenching, as evident from Figure 4.12b.

98 a) b)

Before HT After HT Before HT After HT

0.35 1.E+01 0.30 1.E+00 1.E-01 0.25 1.E-02 1.E-03 0.20 1.E-04 1.E-05 0.15 Current (A) Current 1.E-06

Current (A) Current 1.E-07 0.10 1.E-08 1.E-09 0.05 1.E-10

0.00 9 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 8 Voltage (V) 7 6 Before heat treatment: After heat treatment: 5 Voc= 632 mV Voc= 600 mV 4 2 2 Jsc= 36 mA/cm Jsc= 3.77 mA/cm 3

FF= 0.660 FF= 0.310 ideality Factor Local 2 η= 15.06 η= 0.7 1 0 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 Voltage (V) Figure 4.10: The I-V characteristics of the cell from group QC with electroless Ni as the barrier layer, before and after a 2 hour heat treatment at 200 °C, followed by quenching in ethylene glycol, a) the illuminated I-V curve, b) the corresponding dark I-V and m-V curves.

a) b)

Figure 4.11: High resolution open circuit PL images of the cell from group QC with electroless Ni as the barrier layer, a) before; b) after a 2 hour heat treatment at 200 °C, followed by a quench in ethylene glycol.

99 a) b)

Before HT After HT Before HT After HT

0.35 1.E+01 0.30 1.E+00 1.E-01 0.25 1.E-02 1.E-03 0.20 1.E-04 1.E-05 0.15 Current (A) Current 1.E-06 Current (A) Current 1.E-07 0.10 1.E-08 1.E-09 0.05 1.E-10 6 0.00 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 5 Voltage (V) 4

Before heat treatment: After heat treatment: 3 Voc= 626 mV Voc= 636 mV 2 2 2 Jsc= 38.6 mA/cm Jsc= 36.4 mA/cm

FF= 0.660 FF= 0.583 ideality Factor Local 1 η= 15.97 η= 13.52 0 0.00 0.10 0.20 0.30 0.40 0.50 0.60 Voltage (V) Figure 4.12: The I-V characteristics of the cell from group QC with LIP Ni as the diffusion barrier layer before and after a 2 hour heat treatment at 200 °C, followed by quenching in ethylene glycol, a) the illuminated I-V curve, b) the corresponding dark I-V and m-V curves.

a) b)

Figure 4.13: High resolution open circuit PL images of the cell from group QC with LIP Ni as the barrier layer, a) before; b) after a 2 hour heat treatment at 200 °C, followed by a quench in ethylene glycol.

100 Figure 4.14a-c present a comparison of the m-V hump characteristics between the cells with electroless Ni and LIP Ni as barrier layers, after being exposed to a thermal stress at 200 °C. If the cells were quenched in ethylene glycol, the FWHM of the hump increased considerably for the cells with electroless Ni, whereas it remained unaffected for the cells that had LIP Ni as the barrier layer (Figure 4.14a). Similarly, the peak position of the hump for the cells with electroless Ni was notably shifted to a larger voltage (12.5% increase for 2 hours treatment, 15.9% increase for 2.5 hours treatment) whereas the cells with LIP Ni exhibited a smaller increase in the repositioning of the m-V hump (3.8% increase for 2 hours treatment, 4.6% increase for 2.5 hours treatment), as shown in Figure 4.14b. These observations suggest that the cells which use a LIP Ni barrier layer in their metallisation process are more resistant to the thermal stress of 200 °C for the case of quenching in ethylene glycol.

Furthermore, the degradation of lifetime after thermal stress can be observed by comparing the PL images of the cells before and after heat treatment. Figure 4‎ .11 and Figure 4‎ .13 present the high resolution open circuit PL images of the cells with electroless Ni and LIP Ni barrier layers from group QC, respectively. Regions of relatively high lifetime appear bright in PL images whereas the dark areas correspond to regions of low lifetime possibly due to creation of recombination centres or shunts in the cells. Significant darkening of the PL image of the cell with electroless Ni barrier layer after heat treatment, particularly around the metal fingers and the surface defects suggest areas of increased shunting.

101 a) d)

Electroless Ni/LIP Cu LIP Ni/LIP Cu Electroless Ni/LIP Cu LIP Ni/LIP Cu

32 32 28 28 24 24 20 20 16 16 12 12 8 8 4 4 0 0 -4 -4 Increase in FWHM (%) FWHM in Increase Increase in FWHM (%) FWHM in Increase -8 -8 Group QC Group QD Group SA Group SB

2:00 hrs at 200 C 2:30 hrs at 200 C 2:00 hrs at 200 C 2:30 hrs at 200 C Quenched Quenched Slow Cooled Slow Cooled

b) e)

Electroless Ni/LIP Cu LIP Ni/LIP Cu Electroless Ni/LIP Cu LIP Ni/LIP Cu

16 16 14 14 12 12 10 10 8 8 6 6 4 4 2 2 0 0 -2 -2 peak shiftpeakto larger voltage (%) peak shiftpeakto larger voltage (%) Group QC Group QD Group SA Group SB

2:00 hrs at 200 C 2:30 hrs at 200 C 2:00 hrs at 200 C 2:30 hrs at 200 C Quenched Quenched Slow Cooled Slow Cooled

c) f)

Electroless Ni/LIP Cu LIP Ni/LIP Cu Electroless Ni/LIP Cu LIP Ni/LIP Cu

54 54 50 50 46 46 42 42 38 38 34 34 30 30 26 26 22 22 18 18 14 14 10 10 6 6 Increase in peakIncreaseheight in (%) 2 peakIncreaseheight in (%) 2 -2 -2 Group QC Group QD Group SA Group SB

2:00 hrs at 200 C 2:30 hrs at 200 C 2:00 hrs at 200 C 2:30 hrs at 200 C Quenched Quenched Slow Cooled Slow Cooled Figure 4.14: A comparison of the m-V hump characteristics between the cells with electroless Ni and LIP Ni as barrier layers, after heat treatments at 200 °C, a-c) the cells were quenched in ethylene glycol, d-f) the cells were cooled slowly under ambient laboratory conditions. The percentage variations in the hump characteristics for the two methods of cooling are presented on the same scale.

102 The LDSE cells used in the experimental work of this thesis exhibited a different behaviour when the cells were allowed to cool down slowly under ambient laboratory conditions after the thermal stress. The analysis of the I-V and m-V curves of the cells from groups SA and SB which were slow-cooled revealed similar results before and after heat treatments. As can be seen from Figure 4.14d-f, the m-V hump of the cells with both electroless Ni and LIP Ni as barrier layers, were hardly affected after the heat treatments at 200 °C. In the worst case, the FWHM of the hump increased by only 5% while the position and the height of the hump peak remained virtually unaffected.

By comparing plots a-c with plots d-f in Figure 4.14, respectively, it can be deduced that the type of the Ni barrier layer does not influence the electrical properties of the cells after a slow cool from the thermal stress at 200 °C. In addition, both cells with electroless Ni and LIP Ni barrier layers, exhibited reasonable tolerance against the thermal stress of 200 °C. The effect of the 200 °C heat treatment following a slow cool on the efficiency of the cells with both electroless Ni and LIP Ni diffusion barriers is illustrated in Figure 4.15. The cell efficiency in the worst case reduced slightly from 16.13% to 15.22%.

18.00 17.00 LIP Ni/LIP Cu Before HT 16.00 15.00 14.00 13.00 LIP Ni/LIP Cu 12.00 After HT 11.00 10.00 9.00 8.00 Electroless Ni/LIP Cu 7.00 Before HT Efficiency (%) Efficiency 6.00 5.00 4.00 Electroless Ni/LIP Cu 3.00 After HT 2.00 1.00 0.00 Group SA Group SB

2:00 hrs at 200 C Slow Cooled 2:30 hrs at 200 C Slow Cooled

Figure 4.15: The effect of the 2 hour heat treatment followed by a slow cool under ambient laboratory conditions on the absolute efficiency of LDSE solar cells. The cells with electroless Ni and LIP Ni as the diffusion barrier layer are marked by green and orange, respectively.

103 Examples of the I-V and m-V curves of representative cells from group SA, before and after the 2 hour thermal stress at 200 °C, followed by a slow cool, are shown in Figure 4.16 and Figure 4.18, with electroless Ni and LIP Ni barrier layers, respectively. No significant difference was observed between the two types of Ni barrier layers for either heat treatment following a slow cool. This observation is in good agreement with the high resolution open circuit PL images of the cells before and after heat treatment shown in Figure 4.17 and Figure 4.19.

104 a) b)

Before HT After HT Before HT After HT

0.30 1.E+00 1.E-01 0.25 1.E-02

0.20 1.E-03 1.E-04 0.15 1.E-05 Current (A) Current 1.E-06 Current (A) Current 0.10 1.E-07 1.E-08 0.05 1.E-09

0.00 5 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70

Voltage (V) 4

Before heat treatment: After heat treatment: 3

Voc= 622 mV Voc= 613 mV 2 2 2 Jsc= 36.6 mA/cm Jsc= 36.2 mA/cm

FF= 0.714 FF= 0.722 ideality Factor Local 1 η= 16.29 η= 16.02 0 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 Voltage (V) Figure 4.16: The I-V characteristics of the cell from group SA with electroless Ni as the barrier layer, before and after a 2 hour heat treatment at 200 °C followed by a slow cool under ambient laboratory conditions, a) the illuminated I-V curve, b) the corresponding dark I-V and m-V curves. a) b)

Figure 4.17: High resolution open circuit PL images of the cell from group SA with electroless Ni as the barrier layer, a) before; b) after a 2 hour heat treatment at 200 °C followed a slow cool under ambient laboratory conditions. The dark square shape area in the top left region is the silver tab on the back of the solar cell.

105 a) b)

Before HT After HT Before HT After HT

0.30 1.E+00 1.E-01 0.25 1.E-02

0.20 1.E-03 1.E-04 0.15 1.E-05 Current (A) Current 1.E-06 Current (A) Current 0.10 1.E-07 1.E-08 0.05 1.E-09 5 0.00 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 4 Voltage (V) 3 Before heat treatment: After heat treatment: Voc= 624 mV Voc= 623 mV 2 2 2 Jsc= 37.7 mA/cm Jsc= 37.8 mA/cm

FF= 0.757 FF= 0.754 ideality Factor Local 1 η= 17.82 η= 17.78 0 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 Voltage (V) Figure 4.18: The I-V characteristics of the cell from group SA with LIP Ni as the barrier layer, before and after a 2 hour heat treatment at 200 °C followed by a slow cool under ambient laboratory conditions, a) the illuminated I-V curve, b) the corresponding dark I-V and m-V curves.

a) b)

Figure 4.19: High resolution open circuit PL images of the cell from group SA with LIP Ni as the barrier layer, a) before; b) after a 2 hour heat treatment at 200 °C following a slow cool under ambient laboratory conditions. The dark square shape area in the top left region is the silver tab on the back of the solar cell.

106 4.4.2 Quantitative Current Measurements using Photoluminescence Imaging

A PL imaging system with a zoom lens was used to obtain open circuit PL images of the LDSE solar cells before and after heat treatment. Based on the method described in Section 4.2.2.1‎ , PL images were then analysed to estimate the changes in the current flow through local shunts in the cells upon thermal stress. The effect of the cooling rate upon termination of thermal treatment on the local recombination centres or shunts and the increase in the current flow through the shunts was also investigated. Analysis was done on two different regions identified with low PL signal (Sections A and B), for the cells in each group.

Figure 4.20a and Figure 4.20b present PL images of the cell with electroless Ni from group QD, before and after a 2.5 hour heat treatment at 200 °C, followed by a quench in ethylene glycol, respectively. The selected analysis areas featuring a low PL signal are displayed in both images. As described in Section 4.2.2.1‎ , the dark areas in the PL image indicate local shunted regions. It is evident from the PL images that the thermal treatment followed by quenching created significant shunting in the cell. Figure 4.20c presents the normalised PL counts in each column of pixels as a function of the horizontal distance through the selected area (Section A), before and after heat treatment. Note the horizontal distance of the selected analysis area for all quantitative shunt calculations is equal to the metal finger spacing which is 1 mm. It is clear that the region of the degraded PL signal became remarkably wider on each side of the metal finger after heat treatment. As a result, the current flowing into the shunt increased significantly from 5.10 × 10-5 A before heat treatment to 1.28 × 10-4 A after heat treatment, corresponding to a relative increase of about 150%.

On the contrary, the current flowing through the shunt in the cell with LIP Ni barrier layer from group QD increased slightly by 18.75% after heat treatment. Figure 4.21 presents the PL images of the cell before and after heat treatment along with the plot profile of one of the selected analysis areas (Section B). As can be seen, shunts in the cell with LIP Ni barrier were not significantly exacerbated as a result of the thermal treatment followed by quenching.

107 a) b)

Section B Section B

Section A Section A 1 mm

0.97

0.94

0.91

0.88

0.85

0.82

0.79

0.76 Normalised PL Counts PL Normalised

0.73 -500 -400 -300 -200 -100 0 100 200 300 400 500 Distance (µm)

Before Heat Treatment After Heat Treatment

Figure 4.20: High resolution open circuit PL images of the cell from group QD with electroless Ni as the barrier layer, showing both selected analysis areas, a) before; b) after a 2.5 hour heat treatment at 200 °C, followed by quenching in ethylene glycol, c) plot profile of the selected analysis section A. The dark rectangular area in the top left region is the silver tab on the back of the solar cell.

108 a) b)

Section B Section B

Section A Section A

1 mm

0.97

0.94

0.91

0.88

0.85

0.82

0.79

Normalised PL Counts PL Normalised 0.76

0.73 -500 -400 -300 -200 -100 0 100 200 300 400 500 Distance (µm)

Before Heat Treatment After Heat Treatment

Figure 4.21: High resolution open circuit PL images of the cell from group QD with LIP Ni as the barrier layer, showing both selected analysis areas, a) before; b) after a 2.5 hour heat treatment at 200 °C followed by quenching in ethylene glycol, c) plot profile of the selected analysis section B.

A comparison of the percentage increase in shunt current for the cells with electroless Ni and LIP Ni barrier layers for the various heat treatments is presented in Figure 4.22. It can be seen that all cells with an electroless Ni barrier layer were more susceptible to increased shunting after the thermal stress of 200 °C for the case of quenching in ethylene glycol. In contrast, the changes in shunt current for the cells that were cooled slowly under ambient laboratory conditions were comparatively minor and apparently independent of the Ni barrier layer type.

The results from the PL imaging analysis are consistent with the results from the I-V measurements presented in Section ‎4.4.1. If the cells were quenched in ethylene glycol after heat treatment at 200 °C, the cells with LIP Ni barrier layer outperformed the cells

109 with electroless Ni as the barrier layer. However, for the case of slow cooling both cell types exhibited similar behaviour and were not severely affected by the heat treatment.

165.00

150.00 Electroless Ni/LIP Cu Section A 135.00

120.00 Electroless Ni/LIP Cu 105.00 Section B (%) 90.00 Rsh 75.00 LIP Ni/LIP Cu 60.00 Section A

45.00 Increase in I in Increase 30.00 LIP Ni/LIP Cu Section B 15.00

0.00 Group QB Group QC Group QD Group SA Group SB

1:30 hrs at 200 C 2:00 hrs at 200 C 2:30 hrs at 200 C 2:00 hrs at 200 C 2:30 hrs at 200 C Quenched Quenched Quenched Slow Cooled Slow Cooled

Figure 4.22: A comparison of the percentage increase in the current flowing through the shunt for the cells with electroless Ni and LIP Ni barrier layers, calculated for two different areas as Sections A and B on each cell. The cells in Groups QB-QD were quenched in ethylene glycol after the heat treatments and the cells in groups SA and SB were slow cooled under ambient laboratory conditions.

4.5 Discussion

4.5.1 The Influence of Plating Method

As described in Section 4.4.1‎ , an m-V hump is typically exhibited by LDSE solar cells (see Figure 4.3 as an example). The hump which is an increase in the local ideality factor to above unity between low to moderate voltages, can potentially reduce the FF and in some cases causes a significant decrease in the Voc. McIntosh proposed that the m-V hump is associated with a region of resistance-limited enhanced recombination, i.e., a region in a solar cell with a high recombination rate which is isolated from the remainder of the cell by a resistive path (McIntosh 2001). A possible region of such recombination in solar cells is at the localised Schottky contacts between the metal and the p-type base (Cotter et al. 2000). In LDSE solar cells, a Schottky contact may form if

110 the metal penetrates through the lightly-doped emitter and touches the underlying p-type Si in localised regions. These Schottky contacts are mainly caused by the laser-induced defects mostly in the form of dislocations and holes on the edges of the laser-doped regions (Sugianto et al. 2007, Tjahjono et al. 2007a, Hameiri 2010a).

The holes in the dielectric layer at the edges of the laser-doped regions are possibly caused by the Gaussian-shaped irradiance distribution of the laser beam. The centre of the beam has the highest energy which is sufficient to melt both the dielectric layer and the underlying Si. The energy of the laser beam is reduced further away from the centre. This energy is sufficient to melt the Si in the regions adjacent to the main laser-doped area but not the overlying dielectric layer which typically has a higher melting point than that of Si. As the Si is melting, thermal stresses build up in the localised regions of the dielectric layer. This causes parts of the dielectric layer to break off or puncture, as illustrated in Figure 4.23. Consequently, the underlying molten Si may escape, leaving behind holes on the Si surface adjacent to the laser-doped area (Sugianto et al. 2007).

Figure 4.23 A cross-sectional schematic diagram of the exposed lightly-doped emitter on the perimeter of the laser-doped line (Used with the permission of Tjahjono 2010).

Furthermore, the thermal expansion coefficient mismatch between the dielectric layer and the underlying Si may cause dislocations between the laser-doped and non laser- doped regions. The thermal expansion of the molten Si may become large enough to damage the overlying dielectric layer. As the dielectric layer breaks away it may simultaneously remove small pieces of Si which causes crystallographic defects on the Si surface in the form of dislocations. Both types of laser-induced defects can expose the underlying lightly-doped emitter or even the p-type bulk in severe cases (Sugianto et al. 2007).

111 A SEM image showing the laser-induced defects on the edges of the laser-doped region is shown in Figure 4.24. These defects are revealed after the Yang defect etch. The Yang etch selectively etches the Si surfaces where there are defects at significantly faster rate than areas with no defects (Yang 1984).

In addition to the abovementioned defects generated by the laser doping process, manual handling of the cells can cause damage to the cell surface, creating cracks or micro-cracks. This may result in breakage of texturisation pyramid tips, exposing the conductive area underneath.

Figure 4.24: Laser-induced defects on the edges of a laser-doped region revealed after the Yang defect etch.

Depending on the type of plating method used, these defects may be plated with metal. In the LIP process, metal plating occurs only on the n-type surface where electrons are generated. This is beneficial to LDSE solar cells where the p-type regions may be exposed due to the laser-induced defects. On the contrary, in the electroless plating technique, the metal is plated on any exposed Si surface, regardless of whether it is n- or p-type. Therefore, the metal is plated on the defects generated during the laser doping process.

112 The plated metal may be driven in deeper into the underlying lightly-doped emitter after the subsequent metal sintering step. As a result, these defects turn into high-recombination regions, but are limited by the high resistance of the lightly-doped emitter which is ~ 85-100 Ω/□ in these experiments. Furthermore, since the exposed lightly-doped emitter has a shallow junction of only 0.2-0.5 µm, the plated metal can penetrate through the junction and therefore contact the p-type base, causing localised shunting and/or Schottky contacts (Sugianto et al. 2007, Tjahjono et al. 2007a). These shunts are also resistance-limited because the p-type bulk is lightly-doped to ~ 1.5 × 1016 cm-3.

The formation of a Schottky diode contact has a significant impact on the electrical performance of the cells. The cells with electroless Ni (Figure 4.10), and LIP Ni (Figure 4.12) exhibit distinctive humps in their local ideality factor curves both before and after heat treatments. The humps clearly observed in these graphs indicate junction recombination and/or localised shunting due to the Schottky diode contact formation. However, the cell with electroless Ni as a barrier layer suffered from significant performance loss after the thermal treatment. As seen in Figure 4.10, the local ideality factor increased from 6.0 to 9.2 and extended from 0.4 V to 0.51 V. This corresponds to the reduction of the FF from 66% to 31% and significant degradation of efficiency from 15% to 0.7%. The significant variation in the electrical performance of the cell with electroless Ni after heat treatment can be explained by the fact that the exposed p-type regions via laser-induced defects are plated with Ni during electroless Ni plating. This already creates damage by forming Schottky diode contacts. In addition, during the subsequent Ni sintering process, the Ni plated on the other defected regions may drive in deeper in the lightly-doped emitter in the vicinity of the junction; thus creating extended junction recombination centres. Or, it can even penetrate through the junction because of the shallow junction depth in that area and therefore create localised, non- linear shunting.

Importantly, the Ni-plated regions are highly conductive; therefore, in the subsequent Cu plating via a LIP process, Cu is plated on these regions. A possible explanation for the observed adverse effects of thermal treatment on the electrical properties of the cell is that the plated Cu on the laser-induced defects on the edges of laser-doped lines penetrates through the Ni layer into the lightly-doped emitter or even through the

113 shallow junction. This can severely degrade the cell performance by causing high junction recombination and shunting, respectively.

In contrast, the electrical performance of the cells with LIP Ni as the barrier layer was less affected by the thermal stress as demonstrated in Figure 4.12. The local ideality factor increased from 5.0 to 5.8 and was only extended from 0.43 V to 0.45 V, which caused the FF to reduce from 66% to 58.3%. The efficiency of the cells slightly reduced from 15.9% to 13.5%. These results indicate lower junction recombination and/or lower Schottky diode contact formation compared to the case of electroless Ni as the barrier layer. Despite the possible exposure of p-type regions due to laser-induced defects, neither Ni nor Cu is plated on these regions via the LIP technique. This reduces the area of the Schottky diode contact. Therefore, the subsequent heat treatment does not cause destructive effects on the electrical properties of the cells.

4.5.2 The Influence of Cooling Conditions

The cooling rate upon termination of heat treatment can significantly affect the electrical properties of the cells. If slowly cooled under ambient laboratory conditions, the electrical properties of the cells were only slightly affected by the thermal treatment. However, for the case of a fast cooling rate by quenching in ethylene glycol, the cells with electroless Ni as the barrier layer were severely degraded after the thermal stress. An explanation for this effect may be related to the process of Cu precipitation in Si. Copper precipitates are active recombination centres which can significantly reduce minority carrier lifetime (Istratov and Weber 1998, Istratov and Weber 2002).

According to previous reports, Cu precipitates in Si during or immediately after cooling down after a high temperature treatment (Istratov and Weber 1998, Seibt et al. 1998). The structure, density and spatial distribution of Cu precipitates in Si depend on the cooling rate, initial Cu concentration and pre-existing crystal defects such as dislocations, grain boundaries and microscopic lattice defects in the wafer (Istratov and Weber 1998, Seibt et al. 1999). In this work the wafer type, sample preparation steps and the heat treatment parameters were identical for each group of cells. Therefore the kinetics of Cu precipitation are examined here to explain the observed dependence of cell electrical performance on the cooling rate after heat treatment.

114 In order for an impurity to precipitate, the chemical driving force for precipitation must overcome the nucleation barrier for precipitation. The driving force for precipitation which is usually related to the supersaturation of the dissolved impurity can be calculated from (Seibt et al. 1998, Istratov and Weber 2002):

퐶 휇 = 푘퐵푇 × 푙푛 (2.39) 퐶0 푇

where C is the concentration of the dissolved metal, C0(T) is the equilibrium solubility of the metal, T is the temperature and kB is the Boltzmann constant. If the sample is rapidly cooled from a high temperature by quenching, C0(T) will rapidly drop with decreasing temperature while there is still a high level of dissolved Cu in Si. This causes a large supersaturation of Cu which drives µ to exceed the nucleation barrier. This, as a result, facilitates formation of new Cu precipitates. The Cu silicide precipitates which are most likely in the form of Cu3Si, are detrimental to the electrical properties of the cells since they form a defect band close to the middle of the Si bandgap and are active recombination centres for minority carriers. Furthermore, it is reported that the nucleation barrier in n-type Si may be decreased because of electrostatic effects which makes Cu precipitation in n-type Si possible at lower supersaturation levels, as opposed to p-type Si. This effect makes the Cu precipitates in n-type Si stable at room temperature (Istratov and Weber 2002). The combination of these properties may be used to explain both the observation of degraded electrical performance of the cells with electroless Ni barrier layer after a fast quench following the heat treatments in groups QA-QD.

On the contrary, a slow cooling rate causes C0(T) to decrease gradually. As a result, the Cu supersaturation level and therefore, µ rises slowly (Istratov and Weber 2002). In this case, Cu will first form precipitates on dislocations or microscopic lattice defects which have smaller nucleation barriers. The precipitation of Cu at these sites causes the dissolved Cu concentration (C) to gradually reduce. Consequently, µ will decrease below the nucleation barrier and therefore fewer new Cu precipitates will form. In this scenario conditions are favourable for growth of additional Cu precipitates on pre- existing defects. The already formed precipitates will therefore enlarge. However, if C does not decrease consistently with the decreasing C0(T), as the sample is cooling, µ

115 will rise again and if it overcomes the nucleation barrier, new Cu precipitates will form. Therefore, a fast cooling rate results in a high density of small Cu precipitates in Si bulk whereas slow cooling rates lead to formation of a low density of large Cu precipitates (Istratov and Weber 2002). In fact, it was reported that the minority carrier diffusion length decreases as the size of Cu precipitates decreases. The formation of high density of small diameter Cu precipitates causes the diffusion length and therefore the lifetime to reduce (Istratov and Weber 1998).

From the observation that the electrical performance of the cells is not significantly affected after a slow cool following heat treatment (groups SA and SB), it may be presumed that the Cu tends to precipitate on defect sites which are already recombination centres, thus precipitate formation may not significantly increase the number of recombination centres per unit volume.

Another possible explanation is that under slow cooling conditions, the diffused Cu may not precipitate, but instead remain interstitially in Si. As evident from Equation 4.34, µ falls with reductions in Cu concentration. At low Cu concentrations fewer precipitates form, especially if the cooling rate is slow. It may be the case that for the experiments of groups SA and SB in this work, where relatively low concentrations of Cu of around 108 cm-3 were diffused at 200 °C, µ did not become sufficiently large to induce precipitate formation when the cells were cooled slowly. Furthermore, high quality CZ wafers were used in these experiments which do not typically contain a high number of crystal defects where Cu precipitates form favourably. This is an additional reason why Cu may tend to remain at interstitial sites in Si instead of precipitating. It has been reported that interstitially dissolved Cu in Si is less detrimental to the minority carrier lifetime than Cu precipitates (Istratov and Weber 1998). This may possibly explain the superior electrical performance of the cells when a slow cooling rate is used.

Another Cu-migration mechanism studied by Shabani et al., is the out-diffusion of Cu to the wafer surfaces in B-doped p-type Si at room temperature (Shabani et al. 1996). The reason may be related to the low thermal stability of CuB pairs. However, Cu out- diffusion in n-type Si was only observed if the wafer was heated up to 400 °C. Possibly this is because Cu may form precipitates or complexes with P in n-type Si which are stable at room temperature (Shabani et al. 1996, Istratov and Weber 2002). As discussed 116 in the TEM/EDX analysis in Chapter 3, most of the Cu precipitates were observed in the heavily-doped (n++) area of the experimental cells of this thesis. To verify the effect of a thermal treatment on Cu out-diffusion from Si, a cell from group QC with electroless Ni as a barrier layer was chosen to undergo a second heat treatment at 250 °C for 2 hours and a subsequent slow cool. The high resolution PL images of the cell before and after each heat treatment are displayed in Figure 4.25. It can be seen from Figure 4.25c that both the overall area of the cell and the dark regions in the vicinity of the metal fingers which were possibly due to severe shunting, appear much brighter after the second heat treatment and slow cooling. Note that the illumination exposure time of the cell PL image after the second heat treatment was half the exposure time of the images for before and after the first heat treatment. Therefore, the observed increase in the PL intensity after the second heat treatment indicates a significant improvement in the lifetime. Shabani et al., reported that CuP pairs may be broken at temperatures higher than 250 °C (Shabani et al. 1996) thereby releasing the trapped Cu. Another possible explanation of this effect is the dissolution of the Cu precipitates by additional thermal treatment. Furthermore, the slow cooling rate may have provided favourable conditions for the released Cu to diffuse out to the surface and other existing crystal defects. Another potential contributing factor to the improved PL is passivation of the p-type Si region by Cu+ ions (Lee and Morrison 1988, Naito and Nakashizu 1992, Istratov and Weber 1998).

The illuminated and dark I-V and m-V curves of the cell were also measured after the second heat treatment and subsequent slow cool (Figure 4.26). In contrast to the lifetime improvement as indicated by the increase in the PL intensity, the electrical performance of the cell did not improve. This may be due to damage to the front metal contacts following the heat treatment. The front metal fingers in LDSE solar cells must be plated to a sufficient thickness so as to reduce the metal resistance and minimize the cell series resistance (Tjahjono et al. 2010 ). The thick metal layer may potentially increase the risk of poor adhesion of the plated metal to the Si surface, particularly with varying temperature in solar cells. This is due to the thermal expansion mismatch between the Si and the metal. Therefore, long exposures to high temperatures can increase the series resistance of the cell which can reduce the device FF and efficiency. In the worst-case scenario, the metal lines may peel away from the Si substrate. This was indeed observed in some of the experiments where the metal lines peeled off while being dried with a N2 117 gun after the quench from the high temperature treatment. Although PL imaging with current extraction can be used to determine the series resistance, it was not possible to perform this measurement due to the incompatible size of the cell with the PL system. However, it is recommended to perform a future study using different methods such as luminescence imaging for the measurement of the spatially resolved series resistance.

Whilst damage to the front metal contacts is a possible reason for the non-recovery of cell efficiency following the second thermal treatment, further investigation is needed to test this hypothesis. Specifically, attention must be given to how the cooling rate and the Ni barrier layer (electroless or LIP) influence potential thermally-induced damage to the front metal contacts.

a) b) c)

Figure 4.25: High resolution PL images of the cell from group QC with electroless Ni as the barrier layer, a) before the first heat treatment (2 hours at 200 °C, followed by quenching in ethylene glycol), under 8 seconds of illumination exposure, b) after the first heat treatment, under 8 seconds of illumination exposure and c) after a second heat treatment (2 hours at 250 °C, followed by slow cooling at room temperature), under 4 seconds of illumination exposure.

118 a) b)

Before HT After HT After second HT Before HT After HT After second HT

0.35 1.E+01 0.30 1.E+00 1.E-01 0.25 1.E-02 1.E-03 0.20 1.E-04 1.E-05 0.15 Current (A) Current 1.E-06 Current (A) Current 1.E-07 0.10 1.E-08 1.E-09 0.05 1.E-10

0.00 9 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 8 Voltage (V) 7 6 Before heat treatment: After heat treatment: 5 V = 632 mV V = 600 mV oc oc 4 J = 36 mA/cm2 J = 3.77 mA/cm2 sc sc 3 FF= 0.660 FF= 0.310

Local ideality ideality Factor Local 2 η= 15.06 η= 0.7 1 0 nd After 2 heat treatment: 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 Voc= 605 mV 2 Voltage (V) Jsc= 3.75 mA/cm FF= 0.295 η= 0.67 Figure 4.26: The I-V characteristics of the cell from group QC with electroless Ni as the barrier layer, before and after the first heat treatment (2 hours at 200 °C, followed by quenching in ethylene glycol), and after the second heat treatment (2 hours at 250 °C, followed by slow cooling at room temperature), a) illuminated I-V curves, b) dark IV and m-V curves.

4.6 Conclusion

In this chapter, the effects of Cu contamination on the performance of LDSE solar cells were investigated. It was shown that the type of Ni barrier layer strongly influences the electrical performance of the cells after being quenched in ethylene glycol following the

200 °C thermal treatments. Large reductions in the efficiency, FF and Voc were observed for the cells with electroless Ni barrier layers. Additionally, signs of electrical degradation were observed in the m-V hump characteristics and PL images of these cells.

Quantitative shunt current measurements by PL imaging revealed an increase in the severity of shunts in the cells with electroless Ni barrier layer after the 200 °C thermal treatments with subsequent quenching in ethylene glycol. The superior performance of the cells with LIP Ni barrier layer under these conditions is attributed to lower junction recombination and/or lower Schottky diode contact formation. In contrast to electroless

119 plating, in the LIP technique, the metal is only plated on the n-type region; therefore, the exposed p-type regions commonly created during the laser doping process are not plated with Ni or Cu.

Current-voltage curve analysis revealed that the cells with both electroless Ni and LIP Ni as barrier layers suffered from a major efficiency degradation after 5 hours of heat treatment at 400 °C, followed by quenching in ethylene glycol. The severe degradation of the performance of both cells made it difficult to establish whether the variations in the I-V characteristics between the two barrier layers were significant under these conditions.

The influence of the cooling method after heat treatment termination on shunt sites and high-recombination regions was investigated. If cells are allowed to slowly cool down under ambient laboratory conditions after a thermal treatment at 200 °C, the electrical properties of the cells are almost independent of the type of the Ni barrier layer, as determined by the I-V analysis and PL images. In addition, similar electrical performance was observed for both cells with electroless Ni and cells with LIP Ni barrier layers following a thermal stress at 200 °C and a subsequent slow cool. Measurements of current flowing into shunt regions showed virtually no variation in the shunt current after this type of heat treatment and m-V characteristics did not change appreciably. The superior electrical performance of the cells when a slow cooling rate is used may be related to the fact that under slow cooling conditions, the Cu supersaturation and therefore, the driving force for Cu precipitation builds up slowly. In this scenario Cu tends to precipitate on pre-existing defect sites which are already recombination centres, furnishing existing defect sites. Cu precipitation of this nature may not significantly increase the number of recombination sites in the sample per unit volume.

It is also possible that under slow cooling conditions, the diffused Cu remains at interstitial sites in Si instead of precipitating, especially in high quality CZ wafers which do not typically contain a high number of favourable precipitation sites for Cu. It is well known that interstitially dissolved Cu in Si is less detrimental to the minority carrier lifetime than Cu precipitates which are active recombination centres.

120 However, the behaviour of Cu precipitation is different in multi-crystalline solar cells which contain numerous dislocations, structural defects and grain boundaries. These defects act as favourable precipitation sites for Cu. Therefore, multi-crystalline solar cells may be more sensitive to Cu; however, further work is needed to evaluate the impact of Cu on the efficiency of multi-crystalline solar cells.

The effect of exposing a cell with electroless Ni barrier layer, which was initially cooled rapidly, to a second heat treatment followed by slow cooling, was also investigated. A significant improvement in the lifetime was observed in PL images which is possibly due to dissociation of CuP pairs in n-type Si and/or dissolution of the Cu precipitates. The slow cooling conditions might also promote Cu out-diffusion to the surface. Conversely, the electrical performance of the cell did not improve which may be attributed to the damage to the front metal contacts caused by the thermal expansion coefficient mismatch between the Si and the metal contacts. However, further investigations are required to assess the impact of thermally-induced damage to the front metal contacts on solar cell efficiency.

121

122 Chapter 5 Conclusions and Future Work

The motivation of this thesis was to evaluate potential problems associated with replacing Ag screen-printed paste with plated Ni and Cu for the front-side metallisation of Si solar cells. Although Cu offers several advantages over screen-printed Ag paste, it is a fast diffuser in Si and consequently it can degrade the performance of solar cell devices. This necessitates the use of a barrier layer such as Ni to prevent Cu diffusion into Si. Nickel barrier layers can be formed in different ways and their diffusion barrier properties may depend on their formation method.

The aim of this thesis was to analyse and compare the Cu diffusion barrier properties of Ni layers, formed by electroless plating and LIP, on single-crystalline Si LDSE solar cells. The study focussed on two distinct analytical approaches. In the first approach, which is described in Chapter 3, the effectiveness of the different Ni barrier layers was compared by chemical identification of Cu regions in the heavily-doped Si under the Ni/Cu metal contacts of heat-treated LDSE cells using EDX spectroscopy in conjunction with TEM. The second approach, which is described in Chapter 4, involved examining the effects of possible Cu contamination on the electrical performance of heat-treated LDSE solar cells. This approach used illuminated and dark I-V measurements to qualitatively examine the sources of high-recombination regions and formation of Schottky diode contacts and PL imaging to quantitatively determine the recombination current in the shunted regions of the cells.

123 5.1 Contributions of this Thesis

The main findings from Chapter 3 are summarised below.

1- Energy dispersive X-ray spectroscopy, in conjunction with TEM, can be used successfully to qualitatively detect Cu which diffuses into the Si substrate of heat- treated LDSE solar cells. Quenching of heat-treated LDSE cells in ethylene glycol was used to ensure fast cooling which reduces the out-diffusion of Cu to the surface during cooling. 2- Copper can penetrate through both electroless-plated and LIP Ni barrier layers and be detected up to ~ 1.8 m from the surface of heavily-doped Si of LDSE solar cells that have been heated for 5 hours at 200 °C or 400 °C and then quenched in ethylene glycol. 3- The Cu-rich particles detected in Si are in the form of closely-packed small precipitates and have similar morphology in cells which had electroless-plated and LIP Ni barriers, although the size of the precipitates is smaller in the electroless-plated cells (~ 10 nm) than in the LIP cells (~ 30-50 nm). 4- In all experiments, the Cu-rich particles were detected within the heavily Phosphorous-doped selective-emitter regions.

The main findings from Chapter 4 are listed below.

1- Significant degradations in cell efficiency, FF and Voc were observed for LDSE cells with electroless-plated Ni barrier layers after heating at 200 °C and subsequent quenching in ethylene glycol (fast cooling). In the worst case, the FF reduced from 66% to 31%, causing severe efficiency degradation from 15% to 0.7%. On the other hand, changes in the electrical performance of cells with LIP Ni barrier layers were comparatively minor. The FF reduced slightly from 66% to 58.3% which caused the efficiency to reduce from 15.9% to 13.5%. This observation was confirmed by quantitative measurements of the increase in the shunt current using PL imaging. The degradation in the electrical performance of the LDSE cells having electroless-plated Ni barrier layers may have been due to Cu plating to Si regions other than the heavily laser-doped regions. When the cells were

124 heated, Cu may have penetrated the Si and been trapped in lightly-doped n-type or p-type regions where it could have more severely impacted the cell performance. 2- The electrical performance of both electroless-plated and LIP Ni barrier layer cells degraded severely after 5 hours of heat treatment at 400 °C and subsequent quenching in ethylene glycol. The electrical performance degraded similarly for all LDSE cells suggesting that at higher temperatures the abovementioned benefits of forming a Ni barrier layer using LIP are no longer evident. 3- Under slow cooling conditions (i.e., no quenching), LDSE cells with electroless- plated and LIP Ni barrier layers exhibited only minor electrical performance degradations after sintering at 200 °C. The extent of the performance degradations was similar for both barrier layer types. Therefore, it was concluded that under slow cooling conditions the electrical properties of the cells are almost independent of the type of Ni barrier layer whereas in the case of fast quenching, the cells with LIP Ni barrier layers outperform the cells with electroless-plated Ni barrier layers. 4- Laser-doped selective-emitter cells with electroless-plated Ni barrier layers that were originally heat-treated followed by fast cooling, were exposed to a second heat treatment at 250 °C for 2 hours, followed by slow cooling. Interestingly, a significant improvement in the lifetime, as evidenced by an increased PL intensity, was observed. Possible explanations for this observation include: (i) the breaking of CuP pairs in n-type Si at temperatures higher than 250 °C; and/or (ii) dissolution of the Cu precipitates by additional thermal treatment resulting in out-diffusion of the released Cu to the surfaces and other existing crystal defects. The electrical performance of these cells did not improve. This may have been due to damage to the front metal contacts (e.g., weakened adhesion caused by the thermal expansion coefficient mismatch between the Si and the metal).

5.2 Future work

The EDX/TEM investigations into diffusion of Cu through electroless-plated and LIP Ni barrier layers into the Si of LDSE solar cells, described in Chapter 3, were conducted only on highly-doped areas under the laser-doped lines of LDSE cells. The formation of Cu precipitates in these regions may not result in significant electrical degradation of cells because the concentration of minority carriers in the heavily-doped regions in the

125 LDSE structure is relatively low. Due to the textured surface of the wafers used in the work reported in this thesis and the limiting requirements of the TEM sample preparation, the formation of Cu precipitates within the lightly-doped emitter regions was not investigated. Since the precipitation of Cu in the lightly-doped emitter layer can significantly degrade the electrical performance of the cells and/or create localised shunting, it is recommended to perform similar EDX/TEM analyses in regions adjacent to the laser-doped regions.

The influence of a second heat treatment on Cu out-diffusion from Si was only briefly investigated during this thesis. Further work is required to evaluate the impact of such thermal treatments on the minority carrier lifetime and the electrical performance of the cell. Furthermore, it is suspected that the irreversible damage to the front metal contacts due to the thermal stress is the main reason for the non-recovery of cell electrical performance following the second thermal treatment. Future investigations are required to examine this hypothesis, with specific attention to the effect of cooling rate and the Ni barrier layer (electroless or LIP). Measurement of the spatially resolved series resistance using PL imaging with current extraction is recommended for further studies. This analysis method could not be performed due to the incompatible size of the cell with the PL system used in this work.

This thesis focused on the effects of Cu diffusion in single-crystalline CZ solar cells which do not contain numerous crystal defects. However, multi-crystalline solar cells contain more favourable precipitation sites for Cu such as dislocations, structural defects and grain boundaries. This may mean that multi-crystalline solar cells are less tolerable to Cu contamination than high quality CZ cells. Therefore, further work is needed to investigate the impact of Cu diffusion on the electrical performance of multi-crystalline solar cells.

5.3 Final Comments

The findings of this thesis suggest that many of commonly-accepted views on the use of Ni barrier layers to prevent Cu diffusion in Si solar cells may need to be challenged and further investigated. Although the experiments reported are initial studies and further work is required to substantiate many of the findings, it does appear that Cu can diffuse through both electroless-plated and LIP Ni barrier layers at temperatures as low as 200 °C. However Cu is extremely mobile in Si and, at least in these studies, appears to precipitate in the heavily-doped Si regions underneath the metal where it causes minimal impact on the

126 electrical performance of LDSE cells. If this is indeed the case more generally for selective- emitter Si solar cells, then issues may exist if Ni/Cu metallisation is used with homogeneous emitter cells.

127

128

References

Antler, M. (1970 ) Gold-Plated Contacts: Effect of Heating on Reliability, Plating 57(6), 615- 618.

Augarten, Y., Trupke, T., Lenio, M., Bauer, J., Breitenstein, O., Weber, J. and Bardos, R. A. (2009) Luminescence Shunt Imaging: Qualitative and Quantitative Shunt Images Using Photoluminescence Imaging, in 24th European Photovoltaic Solar Energy Conference, Hamburg, Germany, 21-25 September 2009, 27-31.

Augarten, Y., Trupke, T., Lenio, M., Bauer, J., Weber, J., Juhl, M. and Breitenstein, O. (submitted January 2011) Calculation of Quantitative Shunt Values Using Photoluminescence Imaging, Progress in Photovoltaics.

Bai, P., Yang, G.-R. and Lu, T.-M. (1990) Intrinsic Cu Gettering at a Thermally Grown Sio[Sub 2]/Si Interface, Journal of Applied Physics, 68(7), 3313-3316.

Bartsch, J., Mondon, A., Bayer, K., Schetter, C., Horteis, M. and Glunz, S. W. (2010) Quick Determination of Copper-Metallization Long-Term Impact on Silicon Solar Cells, Journal of The Electrochemical Society, 157(10), H942-H946.

Baumann, J., Kaufmann, C., Rennau, M., Werner, T. and Gessner, T. (1997) Investigation of Copper Metallization Induced Failure of Diode Structures with and without a Barrier Layer, Microelectronic Engineering, 33(1-4), 283-291.

Brady, T. E. and Hovland, C. T. (1981) Scanning Auger Microprobe Study of Gold--Nickel-- Copper Diffusion in Thin Films, Journal of Vacuum Science and Technology, 18(2), 339-342.

Bruker-AXS (2006) Bruker Advanced X-Ray Solutions, Quantax User Manual, Microanalysis System, Based on an Energy-Dispersive X-Ray Spectrometer (Eds) Berlin, Germany: Bruker-AXS Microanalysis GmbH.

Bruton, T., Mason, N. and Summers, J. G. (1992) Towards Production of High Efficiency Terrestrial Solar Cells, in Proceedings of Proceedings of 6th International Photovoltaic Science and Engineering Conference, New Delhi, India, 10-14 February 1992, pp. 11– 16.

Buonassisi, T., Marcus, M. A., Istratov, A. A., Heuer, M., Ciszek, T. F., Lai, B., Cai, Z. and Weber, E. R. (2005) Analysis of Copper-Rich Precipitates in Silicon: Chemical State, Gettering, and Impact on Multicrystalline Silicon Solar Cell Material, Journal of Applied Physics, 97(6), 063503-063503-9.

Cabral, J. C., Harper, J. M. E., Holloway, K., Smith, D. A. and Schad, R. G. (1992) Preparation of Low Resistivity Cu--1 At. %Cr Thin Films by Magnetron Sputtering, in Proceedings of Seattle, Washington (USA), 10, pp. 1706-1712.

Chalmers, B. (1982) The Structure and Properties of Solids: An Introduction to Materials Science, HEYDEN.

Chunduri, S. K. (2009) Be Selective!, Photon International, November, 2009, pp. 108-116. 129

Cotter, J. E., Mehrvarz, H. R., McIntosh, K. R., Honsberg, C. B. and Wenham, S. R. (2000) Combined Emitter and Groove Diffusion in Buried Contact Solar Cells, in Proceedings of 16th European Photovoltaic Solar Energy Conference, Glasgow.

Das, G. (1973) Precipitation of Copper in Silicon, Journal of Applied Physics, 44(10), 4459- 4467.

Davis, J. R., Rohatgi, A., Hopkins, R. H., Blais, P. D., Rai-Choudhury, P., McCormick, J. R. and Mollenkopf, H. C. (1980) Impurities in Silicon Solar Cells, Electron Devices, IEEE Transactions on, 27(4), 677-687. de Moor, H. H. C., Hoornstra, J., Weeber, A. W., Burgers, A. R. and Sinke, W. C. (1997) Printing High and Fine Metal Lines Using Stencils, in 14th EPSEC, Barcelona, 404- 407.

Deng, F., Johnson, R. A., Asbeck, P. M., Lau, S. S., Dubbelday, W. B., Hsiao, T. and Woo, J. (1997) Salicidation Process Using Nisi and Its Device Application, Journal of Applied Physics, 81(12), 8047-8051.

Durkee, L. F. (1979) Method of Plating by Means of Light, US Patent 4144139.

Ebong, A. U., LEE, S. H., Honsberg, C. and Wenham, S. (1996) High Efficiency Double Sided Buried Contact Silicon Solar Cells, Japanese journal of applied physics, 35(4), 2077- 2080.

EnergyMatters (2009) Suntech Pluto Solar Panel 195 Watt 24 Volt (Mono-Crystalline), http://www.energymatters.com.au/suntech-pluto-solar-panel-195-watt-24-volt- monocrystalline-p-2516.html [accessed on 15 November 2011].

Fellmeth, T., Menkoe, M., Clement, F., Biro, D. and Preu, R. (2010) Highly Efficient Industrially Feasible Metal Wrap through (Mwt) Silicon Solar Cells, Solar Energy Materials and Solar Cells, 94(12), 1996-2001.

Fortis Bank Nederland/ VM Group (2010) Silver Book, http://www.virtualmetals.co.uk/pdf/FBNSB0610.pdf [accessed on 2 February 2011].

Fisher, K. (2007) The Pitfalls of Pit Contacts: Electroless Metallization for C-Si Solar Cells, Masters Thesis, The University of New South Wales, Sydney, Australia.

Flink, C., Feick, H., McHugo, S. A., Seifert, W., Hieslmair, H., Heiser, T., Istratov, A. A. and Weber, E. R. (2000) Out-Diffusion and Precipitation of Copper in Silicon: An Electrostatic Model, Physical Review Letters, 85(23), 4900.

Frant, M. S. (1961) Copper Oxides on the Surface of Gold Plate, Plating, 48, 1305-1309.

Gambino, J. P. and Colgan, E. G. (1998) Silicides and Ohmic Contacts, Materials Chemistry and Physics, 52(2), 99-146.

Gauglitz, G. and Vo-Dinh, T., eds. (2003) Handbook of Spectroscopy, Weinheim: Wiley-Vch, [1], page 502.

130 Gee, J. M., Schubert, W. K. and Basore, P. A. (1993) Emitter Wrap-through Solar Cell, in Proceedings of Proceedings of the 23rd IEEE Photovoltaic Specialists Conference, New York, USA, pp. 265-270.

Glunz, S. W. (2007) High-Efficiency Crystalline Silicon Solar Cells, Advances in OptoElectronics, 2007.

Glunz, S. W., Aleman, M., Bartsch, J., Bay, J., Bayer, K., Bergander, R., Filipovic, A., Greil, S., Grohe, A., Hörteis, M., Knorz, A., Menkö, M., Pysch, D., Radtke, V., Rudolph, D., Rublack, T., Schetter, C., Schmidt, D., Woehl, R., Mette, A., Richter, P. and Schultz, O. (2008) Progress in Advanced Metallization Technology at Fraunhofer Ise., in 33rd IEEE Photovoltaic Specialists Conference, San Diego, May 11-16,

Goodhew, P. J., Humphreys, J. and Beanland, R. (2001) Electron Microscopy and Analysis, Third Edition ed., Taylor & Francis.

Graff, K. (1995) Metal Impurities in Silicon-Device Fabrication, Berlin ; New York : Springer- Verlag

Graham, A. H., Lindsay, R. W. and Read, H. J. (1965) The Structure and Mechanical Properties of Electroless Nickel, Journal of the Electrochemical Society, 112(4), 401-413.

Green, M. A. (1995) Silicon Solar Cells : Advanced Principles & Practice, Bridge Printery, Sydney, Australia.

Green, M. A. (1998) Solar Cells: Operating Principles, Technology, and System Applications, Sydney, Australia: The University of New South Wales.

Green, M. A. (2001) 'Crystalline Silicon Solar Cells' in Archer, M. D. and Hill, R., eds., Clean Electricity from Photovoltaics, London: Imperial College Press, http://www.icpress.co.uk/physics/p139.html [accessed on 29 November 2011].

Guo, J.-H. and Cotter, J. E. (2005) Metallization Improvement on Fabrication of Interdigitated Backside and Double Sided Buried Contact Solar Cells, Solar Energy Materials and Solar Cells, 86(4), 485-498.

Guo, J. H. (2004) High Efficiency N-Type Laser-Grooved Buried Contact Silicon Solar Cells, UNSW,

Hameiri, Z. (2010a) Laser-Doped Selective Emitter and Local Back Surface Field Solar Cells with Rear Passivation, The University of New South Wales,

Hameiri, Z. (2010b) Laser-Doped Selective Emitter and Local Back Surface Field Solar Cells with Rear Passivation, Ph.D. Thesis, The University of New South Wales, Sydney, Australia.

Hameiri, Z., Tjahjono, B. S., Mai, L., Wang, S., Sugianto, A., Javid, S. and Wenham, S. R. (2008) Double Sided Laser Doped Solar Cell, in Proceedings of Proceedings of the 18th International Photovoltaic Science and Engineering Conference, Kolkata, India.

Hernandez, J. L., Tous, L., Allebe, C., Philipsen, H., Schlenker, E., John, J., Baert, K. and Poortmans, J. (2010) Application of Cmos Metal Barriers to Copper Plated Silicon Solar Cells, in Proceedings of the 25th European Photovoltaic Solar Energy

131 Conference and Exhibition - EU PVSEC / 5th World Conference on Photovoltaic Energy Valencia, Spain, 1479-1483.

Hilali, M. M., To, B. and Rohatgi, A. (2004) A Review and Understanding of Screen-Printed Contacts and Selective–Emitter Formation, in 14th Workshop on Crystalline Silicon Solar Cells and Modules, Winter Park, Colorado August 8-11, 2004

Hopkins, R. H. and Rohatgi, A. (1986) Impurity Effects in Silicon for High Efficiency Solar Cells, Journal of Crystal Growth, 75(1), 67-79.

Imagej, V 1.44p.(2011) ImageJ, http://imagej.nih.gov/ij/ [accessed on 31 January 2011].

Istratov, A. A., Flink, C., Hieslmair, H., McHugo, S. A. and Weber, E. R. (2000a) Diffusion, Solubility and Gettering of Copper in Silicon, Materials Science and Engineering B, 72(2-3), 99-104.

Istratov, A. A., Flink, C. and Weber, E. R. (2000b) Impact of the Unique Physical Properties of Copper in Silicon on Characterization of Copper Diffusion Barriers, physica status solidi (b), 222(1), 261-277.

Istratov, A. A., Hedemann, H., Seibt, M., Vyvenko, O. F., Schroter, W., Heiser, T., Flink, C., Hieslmair, H. and Weber, E. R. (1998) Electrical and Recombination Properties of Copper-Silicide Precipitates in Silicon, Journal of the Electrochemical Society, 145(11), 3889-3898.

Istratov, A. A. and Weber, E. R. (1998) Electrical Properties and Recombination Activity of Copper, Nickel and Cobalt in Silicon, Applied Physics A: Materials Science & Processing, 66(2), 123-136.

Istratov, A. A. and Weber, E. R. (2002) Physics of Copper in Silicon, Journal of the Electrochemical Society, 149(1), G21-G30.

Itsumi, M., Sato, Y., Imai, K. and Yabumoto, N. (1997) Characterization of Metallic Impurities in Si Using a Recombination-Lifetime Correlation Method, Journal of Applied Physics, 82(7), 3250-3255.

Kasemann, M., Grote, D., Walter, B., Warta, W., Trupke, T., Augarten, Y., Bardos, R. A., Pink, E. and Abbott, M. D. (2007) Shunt Detection Capabilities of Luminescence Imaging on Silicon Solar Cells, in 22nd European Photovoltaic Solar Energy Conference and Exhibition, Milan, Italy,

Kim, D. S., Lee, E. J., Kim, J. and Lee, S. H. (2005) Low-Cost Contact Formation of High- Efficiency Crystalline Silicon Solar Cells by Plating, Journal of The Korean Physical Society, 46(5), 1208-1212.

Kray, D., Aleman, M., Fell, A., Hopman, S., Mayer, K., Mesec, M., Muller, R., Willek, G. P., Glunz, S. W., Bitnar, B., Neuhaus, D.-H., Lüdemann, R., Schlenker, T., Manz, D., Bentzen, A., Sauar, E., Pauchard, A. and Richerzhagen, B. (2008) Laser-Doped Silicon Solar Cells by Laser Chemical Processing (Lcp) Exceeding 20% Efficiency, in 33rd IEEE Photovoltaic Specialist Conference, St. Diego, CA, 12-16 May,

Kray, D., Bay, N., Cimiotti, G., Kleinschmidt, S., Ko, x, sterke, N., Lo, sel, A., Sailer, M., Tra, ger, A., Ku, hnlein, H., Nussbaumer, H., Fleischmann, C. and Granek, F. (2010) Industrial Lcp Selective Emitter Solar Cells with Plated Contacts, in Proceedings of

132 Photovoltaic Specialists Conference (PVSC), 2010 35th IEEE, 20-25 June 2010, pp. 000667-000671.

Kunz, O. (2009) Evaporated Solid-Phase Crystallised Poly-Silicon Thin Film Solar Cells on Glass, The University of New South Wales,

Lauermann, T., Dastgheib-Shirazi, A., Book, F., Raabe, B., Hahn, G., Haverkamp, H., Habermann, D., Demberger, C. and Schmid, C. (2009) Insect: An Inline Selective Emitter Concept with High Efficiencies at Competitive Process Costs Improved with Inkjet Masking Technology, in 24th European Photovoltaic Solar Energy Conference, Hamburg, Germany, September 21-25 2009,

Lee, C. Y., Huang, T. H. and Lu, S. C. (1998) Diffusion Barrier Properties of Electroless Ni for Electroless Cu Using Cu Plating Employing Hypophosphite as a Reducing Agent, Journal of Materials Science-Materials in Electronics, 9(5), 337-346.

Lee, E. J., Kim, D. S. and Lee, S. H. (2002) Ni/Cu Metallization for Low-Cost High-Efficiency Perc Cells, Solar Energy Materials and Solar Cells, 74, 65-70.

Lee, J.-G. and Morrison, S. R. (1988) Copper Passivation of Dislocations in Silicon, Journal of Applied Physics, 64(12), 6679-6683.

Lee, S. H. (2009) Cost Effective Process for High-Efficiency Solar Cells, Solar Energy, 83(8), 1285-1289.

Lennon, A. (2010) Dielectric Patterning for Commercial High-Efficiency Silicon Solar Cells, Ph.D. Thesis, The University of New South Wales, Sydney, Australia.

Liu, D. L., Yang, Z. G. and Zhang, C. (2010) Electroless Ni-Mo-P Diffusion Barriers with Pd- Activated Self-Assembled Monolayer on Sio2, Materials Science and Engineering B- Advanced Functional Solid-State Materials, 166(1), 67-75.

Mai, L., Hameiri, Z., Tjahjono, B., Wenham, S. R., Sugianto, A. and Edwards, M. B. (2009) Rear Junction Laser Doped Solar Cells on Cz N-Type Silicon, in 34th IEEE Photovoltaic Specialists Conference, 1811-1815.

Mallory, G. and Hajdu, J. (1990) Electroless Plating: Fundamentals and Applications, New York: William Andrew Publishing/Noyes Publication.

Mason, N., Russell, R., Artigao, A., Fernandez, J., Nast-Hartley, O., Sherborne, J., Banda, P., Bueno, R., Martinez, G. and Bruton, T. (2004) New Generation BP Solar Saturn Cell in Production: The BP7180 Module, in Proceedings of Proceedings of 19th Symposium PV Solarenergie, Kloster Banz (Staffelstein), 10-12 March 2004.

Materials Evaluation and Engineering, I. (2010) The Handbook of Analytical Methods for Material (Hamm), Materials Evaluation and Engineering, Inc.

McHugo, S. A., Thompson, A. C., Lamble, G., Flink, C. and Weber, E. R. (1999) Metal Impurity Precipitates in Silicon: Chemical State and Stability, Physica B: Condensed Matter, 273-274, 371-374.

McIntosh, K. R. (2001) Humps, Lumps and Bumps: Three Detrimental Effects in the Current- Voltage Curve of Silicon Solar Cells, The University of New South Wales, Sydney.

133 Mette, A., Pysch, D., Emanuel, G., Erath, D., Preu, R. and Glunz, S. W. (2007) Series Resistance Characterization of Industrial Silicon Solar Cells with Screen-Printed Contacts Using Hotmelt Paste, Progress in Photovoltaics: Research and Applications, 15(6), 493-505.

Mette, A., Schetter, C., Wissen, D., Lust, S., Glunz, S. W. and Willeke, G. (2006) Increasing the Efficiency of Screen-Printed Silicon Solar Cells by Light-Induced Silver Plating, in 4th World Conference on Photovoltaic Energy Conversion, Waikoloa, HI 1056-1059.

Miyazaki, M. S., Masakazu; Sumita, Shigeo; Fujino, Nobukatsu (1991) Influence of Metal Impurities on Leakage Current of Si N+P Diode, Japanese Journal of Applied Physics, 30(2B), L295-L297.

Naito, S. and Nakashizu, T. (1992) Defect Engineering in Semiconductor Growth, Processing and Device Technology, in Proceedings of Materials Research Society Symposium Proceedings, 262, pp. 641.

Nes, E. and Lunde, G. (1972) Copper Precipitate Colonies in Silicon, Journal of Applied Physics, 43(4), 1835-1837.

Nes, E. and Washburn, J. (1973) Transmission Electron Microscope Investigation of the Growth of Copper Precipitate Colonies in Silicon, Journal of Applied Physics, 44(8), 3682- 3688.

O'Sullivan, E. J., Schrott, A. G., Paunovic, M., Sambucetti, C. J. and al, e. (1998) Electrolessly Deposited Diffusion Barriers for Microelectronics IBM Journal of Research and Development 42(5), 607-620

Olowolafe, J. O., Nicolet, M. A. and Mayer, J. W. (1976) Influence of the Nature of the Si Substrate on Nickel Silicide Formed from Thin Ni Films, Thin Solid Films, 38(2), 143- 150.

Pinnel, M. R. and Bennett, J. E. (1976) Qualitative Observations on Diffusion of Copper and Gold through a Nickel Barrier, Metallurgical Transactions a-Physical Metallurgy and Materials Science, 7(5), 629-635.

Ralph, E. L. (1975) Recent Advancements in Low Cost Solar Cell Processing, in Proceedings of 11th Photovoltaic Specialists Conference, Scottsdale, Arizona, USA, pp. 315-316.

Riegel, S., Raabe, B., Petres, R., Dixit, S., Zhou, L. and Hahn, G. (2008) Towards Higher Efficiencies for Crystalline Silicon Solar Cells Using Sic Layers, in 23rd European Photovoltaic Solar Energy Conference, Valencia, Spain, September 1-5, 2008,

Rizk, R., Portier, X., Allais, G. and Nouet, G. (1994) Electrical and Structural Studies of Copper and Nickel Precipitates in a Sigma =25 Silicon Bicrystal, Journal of Applied Physics, 76(2), 952-958.

Rohan, J. F. and O'Riordan, G. (2003) Characterisation of the Electroless Nickel Deposit as a Barrier Layer/under Bump Metallurgy on Ic Metallisation, Microelectronic Engineering, 65(1-2), 77-85.

Rohan, J. F., O‟Riordan, G. and Boardman, J. (2002) Selective Electroless Nickel Deposition on Copper as a Final Barrier/Bonding Layer Material for Microelectronics Applications, Applied Surface Science, 185(3-4), 289-297.

134

Schmit, R. R. and Gee, J. M. (2010) Emitter Wrap-through Back Contact Solar Cells on Thin Silicon Wafers,

Seibt, M. and Graff, K. (1988) Characterization of Haze-Forming Precipitates in Silicon, Applied Physics, 63(9), 4444

Seibt, M., Griess, M., Istratov, A. A., Hedemann, H., Sattler, A. and Schröter, W. (1998) Formation and Properties of Copper Silicide Precipitates in Silicon, physica status solidi (a), 166(1), 171-182.

Seibt, M., Hedemann, H., Istratov, A. A., Riedel, F., Sattler, A. and Schröter, W. (1999) Structural and Electrical Properties of Metal Silicide Precipitates in Silicon, physica status solidi (a), 171(1), 301-310.

Shabani, M. B., Yoshimi, T. and Abe, H. (1996) Low-Temperature out-Diffusion of Cu from Silicon Wafers, Journal of the Electrochemical Society, 143(6), 2025-2029.

Shi, Z., Wenham, S. and Ji, J. (2009) Mass Production of the Innovative Pluto Solar Cell Technology, in Proceedings of Photovoltaic Specialists Conference (PVSC), 2009 34th IEEE, 7-12 June 2009, pp. 001922-001926.

Shockley, W. (1949) The Theory of P-N Junctions in Semiconductors and P-N Junction Transistors Bell System Technical Journal, 28, 435-489.

Sinton, R. A., King, R. R. and Swanson, R. M. (1989) Novel Implementations of Backside Contact Silicon Solar Cell Designs in One-Sun and Concentrator Applications, in Proceedings of Conference Record 4th Photovoltaic Science and Engineering, Sydney,, pp. 142.

Solarbuzz (2011) http://www.solarbuzz.com/facts-and-figures [accessed on 30 November 2011].

Solberg, J. (1978) The Crystal Structure of Η-Cu3si Precipitates in Silicon, Acta Crystallographica Section A, 34(5), 684-698.

Sugianto, A., Tjahjono, B. S., Guo, J. H. and Wenham, S. R. (2007) Impact of Laser Induced Defects on the Performance of Solar Cells Using Localised Laser Doped Regions beneath the Metal Contact, in Proceedings of 22nd European Photovoltaic Solar Energy Conference, Milan, Italy, pp. 1759-1762.

Swanson, R. M., Beckwith, S. K., Crane, R. A., Eades, W. D., Young Hoon, K., Sinton, R. A. and Swirhun, S. E. (1984) Point-Contact Silicon Solar Cells, Electron Devices, IEEE Transactions on, 31(5), 661-664.

Szlufcik, J., Sivoththaman, S., Nlis, J. F., Mertens, R. P. and Van Overstraeten, R. (1997) Low- Cost Industrial Technologies of Crystalline Silicon Solar Cells, Proceedings of the IEEE, 85(5), 711-730.

Thompson, A. C., Kirz, J., Attwood, D. T., Gullikson, E. M., Howells, M. R., Kortright, J. B., Yanwei, L., Robinson, A. L., Vaughan, D., Underwood, J. H., Kim, K. J., Lindau, I., Pianetta, P., Winick, H., Williams, G. P. and Scofield, J. H. (2009) X-Ray Data Booklet, Third ed., Lawrence Berkeley National Laboratory, University of California.

135 Tjahjono, B., Wang, S., Sugianto, A., Mai, L., Hameiri, Z., Borojevic, N., Ho-Baillie, A. and Wenham, S. R. (2008) Application of Laser Doped Contact Structure on Multicrystalline Solar Cells, in Proceedings of 23rd European Photovoltaic Solar Energy Conference and Exhibition, Valencia, Spain, September 2008, pp. 1995 - 2000.

Tjahjono, B. S., Guo, J.-H., Hameiri, Z., Mai, L., Sugianto, A., Wang, S. and Wenham, S. R. (2007a) High Efficiency Solar Cell Structures through the Use of Laser Doping, in Proceedings of Proceedings of the 24th European Photovoltaic Solar Energy Conference, Milan, Italy, September 2007.

Tjahjono, B. S., Guo, J. H., Hameiri, Z., Mai, L., Sugianto, A., Wang, S. and Wenham, S. R. (2007b) High Efficiency Solar Cell Structures through the Use of Laser Doping, in Proceedings of 22nd European Photovoltaic Solar Energy Conference, Milan, Italy, pp. 966-969.

Tjahjono, B. S., Yang, M. J., Lan, C. Y., Ting, J., Sugianto, A., Ho, H., Kuepper, N., Beilby, B., Szpitalak, T. and Wenham, S. R. (2010) 18.9% Efficient Laser Doped Selective Emitter Solar Cell on Industrial Grade P-Type Czochralski Wafer, in 25th European Photovoltaic Solar Energy Conference, Valencia, Spain, 1396-1400.

Trupke, T. and Bardos, R. A. (2005a) Photoluminescence: A Surprisingly Sensitive Lifetime Technique, in Proceedings of 31st IEEE Photovoltaic Specialists Conference, pp. 903- 906.

Trupke, T. and Bardos, R. A. (2005b) Photoluminescence: A Versatile Characterisation Technique for Crystalline Silicon, in 15th Workshop on Crystalline Silicon Solar Cells & Modules,

Trupke, T., Bardos, R. A., Abbott, M. D., Chen, F. W., Cotter, J. E. and Lorenz, A. (2006a) Fast Photoluminescence Imaging of Silicon Wafers, in Proceedings of Photovoltaic Energy Conversion, Conference Record of the 2006 IEEE 4th World Conference on, May 2006, 1, pp. 928-931.

Trupke, T., Bardos, R. A., Abbott, M. D., Fisher, K., Bauer, J. and Breitenstein, O. (2006b) Luminescence Imaging for Fast Shunt Localisation in Silicon Solar Cells and Silicon Wafers, in International Workshop on Science and Technology of Crystalline Silicon Solar Cells, Sendai, Japan,

Trupke, T., Bardos, R. A., Schubert, M. C. and Warta, W. (2006c) Photoluminescence Imaging of Silicon Wafers, Applied Physics Letters, 89(4), 044107-044107-3.

Turn, J. C. and Owen, E. L. (1974) Metallic Diffusion Barriers for the Copper- Electrodeposited Gold System, Plating, 61, 1015-18.

Utama, R., Lennon, A., Lenio, M., Borojevic, N., Karpour, A., Ho-Baillie, A. and Wenham, S. (2008) Inkjet Printing for High Efficiency Selective Emitter Silicon Solar Cell, in Proceedings of Proceeding of 23rd European Photovoltaic Solar Energy Conference and Exhibition, Valencia, Spain.

Van Kerschaver, E., Einhaus, R., Szlufcik, J., Nijs, J. and Mertens, R. (1998) A Novel Silicon Solar Cell Structure with Both External Polarity Contacts on the Back Surface, in 2nd World Conference on Photovoltaic Solar Energy Conversion, 1479-1482.

136 Weber, E. R. (1983) Transition Metals in Silicon, Applied Physics A: Materials Science & Processing, 30(1), 1-22.

Wenham, S., Green, M., Watt, M. and Corkish, R. (2006a) Applied Photovoltaics, London, The United Kingdom: Earthscan

Wenham, S., Green, M (1988) Buried Contact Solar Cell, 4726850.

Wenham, S., Shi, Z., Mai, L., Tjahjono, B. and Ji, J. (2006b) Overcoming the Fundamental Performance Limitations of Screen-Printed Solar Cells., in Proceedings of Proceedings of16th Workshop on Crystalline Silicon Solar Cells & Modules: Materials and Processes Denver, Colorado, August 6–9, 2006 pp. 103-106.

Wenham, S. R. (1986) Laser Grooved Silicon Solar Cells, Ph.D. Thesis, The University of New South Wales, Sydney, Australia.

Wenham, S. R. and Green, M. A. (1988a) Buried Contact Solar Cell, Patent 4726850.

Wenham, S. R. and Green, M. A. (1988b) Method of Making Buried Contact Solar Cell, U.S. Patent 4748130.

Wenham, S. R. and Green, M. A. (1996) Silicon Solar Cells, Progress in Photovoltaics Research and Applications, 4, 3-33.

Würfel, P. (1982) The Chemical Potential of Radiation, Journal of Physics C: Solid State Physics, 15, 3967-3985.

Würfel, P. and Ruppel, W. (1981) The Chemical Potential of Luminescent Radiation, Journal of Luminescence, 24-25(Part 2), 925-928.

Würfel, P. and Wü rfel, U. (2009) Physics of Solar Cells : From Basic Principles to Advanced Concepts, Wiley-VCH.

Yang, K. H. (1984) An Etch for Delineation of Defects in Silicon, Journal of The Electrochemical Society, 131(5), 1140-1145.

Yao, Y. (2009) Photoplating for High Efficiency Laser Doped Silicon Solar Cells, Undergraduate Honour Thesis, The University of New South Wales, Sydney, Australia.

Zhao, J., Wang, A., Altermatt, P. P., Wenham, S. R. and Green, M. A. (1996) 24% Efficient Perl Silicon Solar Cell: Recent Improvements in High Efficiency Silicon Cell Research, Solar Energy Materials and Solar Cells, 41-42, 87-99.

Zhao, J., Wang, A. and Green, M. A. (1999) 24.5% Efficiency Silicon Pert Cells on Mcz Substrates and 24.7% Efficiency Perl Cells on Fz Substrates, Progress in Photovoltaics: Research and Applications, 7, 471-474

137