Solar Cells and Light Management: Materials, Strategies and Sustainability 0081027621, 9780081027622

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SOLAR CELLS AND LIGHT MANAGEMENT Materials, Strategies and Sustainability

Edited by

FRANCESCO ENRICHI Research Fellow Department of Molecular Sciences and Nanosystems Ca’ Foscari University of Venice Venice, Italy

GIANCARLO C. RIGHINI Emeritus Research Director Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi Rome, Italy

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States Copyright Ó 2020 Elsevier Ltd. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library ISBN: 978-0-08-102762-2 For information on all Elsevier publications visit our website at https://www.elsevier.com/books-and-journals Publisher: Matthew Deans Acquisition Editor: Gwen Jones Editorial Project Manager: Isabella C. Silva Production Project Manager: R.Vijay Bharath Cover Designer: Christian J. Bilbow Typeset by TNQ Technologies

To my wife Valentina and my children Marco, Davide, and Elena, for their endless love, trust, and support Francesco Enrichi To my Dearest Ones, Marta and Nicoletta Giancarlo C. Righini

Contributors  Hugo Aguas i3N/CENIMAT, Department of Materials Science, Faculty of Science and Technology, Universidade NOVA de Lisboa and CEMOP/UNINOVA, Campus de Caparica, Caparica, Portugal  Mario Enrique Alvarez-Ramos Posgrado en Nanotecnología, Departamento de Física, Universidad de Sonora, Hermosillo, Sonora, México Arturo Ay on MEMS Research Lab, Department of Physics and Astronomy, University of Texas at San Antonio, San Antonio, TX, United States Jérémy Barbé SPECIFIC, College of Engineering, Swansea University, Swansea, United Kingdom Antonino Bartolotta CNR-IPCF, Istituto per i Processi Chimico-Fisici, Messina, Italy Brigitte Boulard IMMM UMR CNRS 6283, Le Mans Université, Le Mans, France Giuseppe Calogero CNR-IPCF, Istituto per i Processi Chimico-Fisici, Messina, Italy Andrea Chiappini IFN-CNR CSMFO Lab. and FBK Photonics Unit, Povo-Trento, Italy Isabella Concina Division of Materials Science, Department of Engineering Sciences and Mathematics, Luleå, Sweden Isodiana Crupi Engineering Department, University of Palermo, Palermo, Italy Paola Delli Veneri Italian National Agency for New Technologies, Energy and Sustainable Economic Development (ENEA), Portici Research Centre, Portici, NA, Italy Aldo Di Carlo Centre for Hybrid and Organic Solar Energy, University of Rome Tor Vergata, Rome, Italy; Laboratory for Advanced Solar Energy, National University of Science and Technology ‘‘MISiS’’, Moscow, Russia Francesco Enrichi Department of Molecular Sciences and Nanosystems, Ca’ Foscari University of Venice, Venice, Italy

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Maurizio Ferrari IFN-CNR CSMFO Lab. and FBK Photonics Unit, Povo-Trento, Italy; Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi, Rome, Italy Alvaro Flores-Pacheco Posgrado en Nanotecnología, Departamento de Física, Universidad de Sonora, Hermosillo, Sonora, México; MEMS Research Lab, Department of Physics and Astronomy, University of Texas at San Antonio, San Antonio, TX, United States Elvira Fortunato i3N/CENIMAT, Department of Materials Science, Faculty of Science and Technology, Universidade NOVA de Lisboa and CEMOP/UNINOVA, Campus de Caparica, Caparica, Portugal Mojtaba Gilzad Kohan Division of Materials Science, Department of Engineering Sciences and Mathematics, Luleå, Sweden Francesco Gonella Department of Molecular Sciences and Nanosystems, Ca’ Foscari University of Venice, Mestre-Venezia, Italy Sirazul Haque i3N/CENIMAT, Department of Materials Science, Faculty of Science and Technology, Universidade NOVA de Lisboa and CEMOP/UNINOVA, Campus de Caparica, Caparica, Portugal Harrison Ka Hin Lee SPECIFIC, College of Engineering, Swansea University, Swansea, United Kingdom Yumiko Katayama Graduate School of Arts and Sciences, The University of Tokyo, Meguro-ku, Tokyo, Japan Anna Lukowiak Institute of Low Temperature and Structure Research, PAS, Wroclaw, Poland Xiaoling Ma Key Laboratory of Luminescence and Optical Information, Ministry of Education, Beijing Jiaotong University, Beijing, China Rodrigo Martins i3N/CENIMAT, Department of Materials Science, Faculty of Science and Technology, Universidade NOVA de Lisboa and CEMOP/UNINOVA, Campus de Caparica, Caparica, Portugal Tiago Mateus i3N/CENIMAT, Department of Materials Science, Faculty of Science and Technology, Universidade NOVA de Lisboa and CEMOP/UNINOVA, Campus de Caparica, Caparica, Portugal Manuel J. Mendes i3N/CENIMAT, Department of Materials Science, Faculty of Science and Technology, Universidade NOVA de Lisboa and CEMOP/UNINOVA, Campus de Caparica, Caparica, Portugal

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Lucia V. Mercaldo Italian National Agency for New Technologies, Energy and Sustainable Economic Development (ENEA), Portici Research Centre, Portici, NA, Italy Itumeleng Mokgosi School of Physics, Centre of Excellence in Strong Materials and Materials for Energy Research Group, University of the Witwatersrand, Johannesburg, South Africa Seweryn Morawiec Institute of Physics, Faculty of Physics, Astronomy and Informatics, Nicolaus Copernicus University, Torun, Poland Dustin R. Mulvaney Department of Environmental Studies, San Jose State University, San Jose, CA, United States Alessandro Lorenzo Palma Centre for Hybrid and Organic Solar Energy, University of Rome Tor Vergata, Rome, Italy; Italian National Agency for New Technologies, Energy and Sustainable Economic Development (ENEA), Energy Efficiency Unit Department, Rome, Italy Alexander Quandt School of Physics, Centre of Excellence in Strong Materials and Materials for Energy Research Group, University of the Witwatersrand, Johannesburg, South Africa Giancarlo C. Righini Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi, Rome, Italy Olalla Sanchez-Sobrado i3N/CENIMAT, Department of Materials Science, Faculty of Science and Technology, Universidade NOVA de Lisboa and CEMOP/UNINOVA, Campus de Caparica, Caparica, Portugal Danila Saranin Laboratory for Advanced Solar Energy, National University of Science and Technology ‘‘MISiS’’, Moscow, Russia Sofia Spagnolo Department of Environmental Sciences, Informatics and Statistics, Ca’ Foscari University of Venice, Mestre-Venezia, Italy Setsuhisa Tanabe Graduate School of Human and Environmental Studies, Kyoto University, Sakyo-ku, Kyoto, Japan Wing Chung Tsoi SPECIFIC, College of Engineering, Swansea University, Swansea, United Kingdom Alberto Vomiero Division of Materials Science, Department of Engineering Sciences and Mathematics, Luleå, Sweden; Department of Molecular Sciences and Nanosystems, Ca’ Foscari University of Venice, Italy

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Robert Warmbier Department of Physics, University of Johannesburg, Auckland Park, South Africa Narges Yaghoobi Nia Centre for Hybrid and Organic Solar Energy, University of Rome Tor Vergata, Rome, Italy Fujun Zhang Key Laboratory of Luminescence and Optical Information, Ministry of Education, Beijing Jiaotong University, Beijing, China Lidia Zur IFN-CNR CSMFO Lab. and FBK Photonics Unit, Povo-Trento, Italy

Preface The growth of global population and the ever-increasing energy consumption put serious questions on the sustainability of our society. In this context, the transition from fossil fuels to renewable energies is recognized as a fundamental strategy to pursue, driven by strong investments and strict regulations. Among renewables, a key role is played by photovoltaics (PV), since sunlight is widespread all over the world and can be exploited at any level, from cheap portable devices up to large industrial installations. In the last few decades, the evolution of solar cells developed incredibly fast, thanks to both a multiplication of different technologies driven by the synthesis and optimization of novel materials and a continuous increase of efficiency, made possible by proper structure design and light management strategies. The aim of this book is to provide the state of the art of the most recent progresses in the field of PV solar cells, with the attempt of shining light on emerging technologies, latest results, and potential perspectives. It is intended both for the expert in the field, and for the newcomer, being a valuable reference tool with updated literature and envisioned developments in the field. It also aims at being an instructive and stimulating text for undergraduate and postgraduate students, allowing a growth of their knowledge about novel materials for solar cells and light management. This book is organized in four sections. After a general introduction, the first section provides a detailed overview on different concepts and materials for solar PV devices, starting from the standard silicon cells to the organicbased, dye-sensitized perovskite and all-oxide solar cells. This panorama allows the reader to get up to date with the standard traditional approach as well as the new and more promising concepts for the evolution of the field. The second and third sections are focused on light manipulation and solar spectrum modification, as tools to further increase the photon-to-current efficiency. This field, too, is in rapid evolution, and the latest results on optical strategies and materials for light trapping, plasmonic scattering, spectral upand downconversion (mainly based on rare earth doped oxides and complexes) are reported. Finally, in the last section, a critical approach on the sustainability of PV technologies is given. An eMergy (spelled with “m”) analysis, which is an extension of the life cycle analysis (LCA), is applied to the study of solar PV systems. Furthermore, the limits of LCA metrics into questions of

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environmental justice, human rights, and other nonquantifiable impacts of production systems are discussed. Nowadays, the attention to issues such as sustainability and social responsibilities cannot be neglected and, instead, should be carefully considered in the development of new technologies and their use. In details, this book starts with an introduction, Chapter 1, which gives an overview of the concepts and a summary of the main results of the PV sector. After a general description of the global scenario and the historical evolution of PV, the different families of solar cells are presented, with their latest efficiency achievements and perspectives of further research development. Furthermore, the main optical strategies for increasing the efficiency of these devices are illustrated with updated results and future potential. Chapter 2 is focused on silicon solar cells from a device-engineering perspective, encompassing both the crystalline and the thin film (based on amorphous and microcrystalline silicon) technologies. Properties and fabrication methods of the photoactive materials are considered, as well as different device architectures and promising new trends for this technology. Organic solar cells (OSCs) are the subject of Chapter 3, with particular attention to the recent advances in ternary structures. The development of ternary OSCs based on different system is summarized; the main working mechanisms and the potential research directions are outlined, together with the challenges and perspectives for future development. An original journey into the field of dye-sensitized solar cells (DSSCs), starting from the colors of chemical dyes to the natural ones, is presented in Chapter 4. The state of the art, design, and fabrication of solar devices based on the appealing family of natural sensitizers are shown, reviewing their lights and shadows. The working principles and the main components of DSSC are described, with a specific focus on the photoanode; a systematic review of the many dyes employed as sensitizers for solar energy application is also provided. In Chapter 5, the emerging field of perovskite solar cells is examined. The chapter summarizes the unique properties of halide perovskite, such as crystal structure, energy band properties, charge transport, recombination, and ferroelectric properties. A detailed discussion is devoted to perovskite crystallization (from liquid phase) and the use of different deposition methods. Furthermore, examples of possible device architectures and module fabrication are discussed. Solar cells based on oxide materials are described in Chapter 6. These oxides, composed of Earth-abundant harmless materials, are more stable than

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most of the other materials alternative to silicon and are therefore ideal candidates for realizing cheap, stable, and safe solar cells. Their major efficiencylimitation factor is the low hole mobility in the p-type oxide. This chapter illustrates the latest results in the field, including the integration of nanowire geometries as a viable solution toward fast charge transport and collection. Chapter 7 illustrates simulation approaches for the modeling of PV materials and devices. The calculation of material parameters by ab initio methods based on density functional theory opens the very exciting possibility to develop new types of solar cells entirely in silico. Furthermore, the chapter gives some insights in the augmentation of commercial types of solar cells by using frequency conversion layers. Metallic nanoparticles sustaining localized surface plasmon resonances are the subject of Chapter 8. The correlation between structural and optical properties of self-assembled silver nanostructures, the optimization of the fabrication process, and their potential for enhancing light trapping in thin film PV are discussed. Furthermore, a novel optoelectronic spectroscopic method for the quantification of absorption enhancement and of parasitic losses in thin photovoltaic absorbers due to plasmonic light trapping is described. Chapter 9 is focused on the development of photonic structures for improving light trapping in thin film solar cells, enabling high-efficiency, low-cost and mechanically flexible devices, in particular for silicon and perovskite PV technologies. In this chapter, colloidal lithography methods are shown to be suitable for surface nanopatterning, allowing compatibility with industrial scalability and low-cost requirements. The field of indoor light harvesting is presented in Chapter 10, in relation to the significant increase of electronic devices for Internet-of-Things (IoT) applications. IoT can be crucial in developing smart homes, offices, and buildings. Compared to silicon, the cheaper organic and perovskite PV cells have much better performance under indoor lighting conditions. The aim of this chapter is to introduce this exciting and emerging application of organic and perovskite PV devices. Chapter 11 discusses transparent glass ceramics activated by rare earth ions as spectral converters to improve the efficiency of solar cells. After a brief presentation of transparent glass ceramics, the basics of spectral conversion are summarized, and some specific examples, mainly based on glass ceramics fabricated by sol-gel route, are reported. In analogy with the previous chapter, Chapter 12 is focused on quantum cutting by lanthanides and Yb3þ co-doped phosphors. After absorption in

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the UV or visible spectral region, two or more photons at about 1 mm wavelength can be emitted, which are better matched to commercial polycrystalline or single crystal silicon absorption. Recent results and updated literature on this field are presented. The use of quantum dots as broadband spectral down-shifters is illustrated in Chapter 13. The experimental observations indicate that higher power conversion efficiency values can be achieved by exploiting this down-shifting effect, due to the absorption of the UV photons by the QDs involved and the subsequent emission at longer wavelengths. In Chapter 14, an eMergy (spelled with “m”) analysis approach is presented, applied to the study of solar PV systems. Emergy analysis is a thermodynamic-based method used for the assessment of systemic sustainability, based on the upstream energy investment necessary for creating a product or a service. This analysis, which is an extension of Life Cycle Assessment (LCA) methodology, allows to determine effective integrated indicators for the evaluation of the global efficiency and the sustainability of the systems at issue. Finally, Chapter 15 draws on several LCA and environmental justice case studies related to PV commodity chains. It aims at improving our understanding of the socio-ecological impacts of global production systems. LCA could be an important tool to explore the roots of environmental inequality, and it brings a way to look at risks in a more comprehensive way, but also poses some challenges. This chapter attempts to provide both the rationale, methodology, and instances of integrating LCA and commodity chain analysis in the assessment of systems of production. This book would not have been possible without the collaboration and the scientific-technical expertise of the authors of the different chapters; we are deeply indebted to them for their excellent and timely work. We are also grateful to our editorial project manager, Isabella Conti Silva, and all the Elsevier editorial staff for the precious and continuous support. Francesco Enrichi Giancarlo C. Righini

CHAPTER ONE

Solar cells’ evolution and perspectives: a short review Giancarlo C. Righini1, Francesco Enrichi2 1

Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi, Rome, Italy Department of Molecular Sciences and Nanosystems, Ca’ Foscari University of Venice, Venice, Italy

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1.1 An introduction: economics, energy, and sustainability The growth in world population and industrialization that started to accelerate in the second half of the 20th century had an obvious effect on the growth of the global economic activities, hence of the world gross product (WGP). In turn, energy consumption has also continuously increased and so did the impact of all these factors on the environment. Doubts were then rising about the possible consequences of continuing to equate growth with progress. The concept of sustainability, “which encompasses the idea of pursuing societal structures and activities that are resilient and enduring, for both human beings and the geobiosphere that supports their life” [1], is relatively recent. A first warning of unlimited material growth and unrestrained consumption in a world which has limited resources came from the work of “Club of Rome,” founded in 1968 [2], and from the publication, in 1972, of the report “The Limits to Growth” [3], that first introduced the idea that environmental and economic issues cannot be treated separately. The idea of assessing the quality of energy necessary to make a product or service and of doing so by taking into account all the relevant parameters (primary energy, economic investment, labor, environmental loading, etc.) led to the introduction of the emergy methodology, largely based on the seminal work by H. T. Odum [4]. An emergy analysis of a sustainable photovoltaic solar exploitation is presented in the Chapter 14 of this book [1]. Fig. 1.1 presents the growth rates of population, WGP, global energy consumption, and electricity consumption in the present decade

Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00001-X

© 2020 Elsevier Ltd. All rights reserved.

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Figure 1.1 Growth rates in the present decade of the population (blue), world gross product (orange), global energy consumption (yellow), and global electricity consumption (gray).

(2010e19). Despite the fact that the population growth slowed in the last years (1.07% growth rate estimated in 2019, with respect to 1.23 in 2010), it reached a count of over 7.7 billions of persons [5]; the GDP annual growth rate fluctuated but was always higher than 3%, reaching 3.6% in 2017 and 2018 (values calculated by adjusting for purchasing power parities, i.e., for differences in the cost of living across countries) [6]. The energy consumption increased at a lower rate than GDP, thanks to a better efficiency of the technologies and industrial processes; in 2017 and 2018; however, there was a faster rise, with a global energy demand increased by 2.1% in 2017, compared with 0.9% the previous year and 0.9% on average over the previous 5 years. More than 40% of the growth in 2017 was driven by China and India. Global energy demand in 2017 reached an estimated 14,050 million tonnes of oil equivalent (Mtoe), compared with 10,035 Mtoe in 2000, according to the “Global Energy and CO2 Status Report 2017,” released by the International Energy Agency (IEA) in March 2018 [7]. The global trend of consumption of electricity, in particular, after a continuous strong increase from 2000 to 2010, slowed down, with a minimum growth rate in 2015, corresponding to minimum growth rate of GDP [8]; then, started again to rise, with an increase of 2.6% in 2017. Global power generation is forecast to increase by some 60% between 2017 and 2040 to cover a quarter of primary energy demand, the IEA said in its World Energy Outlook 2018 [9].

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Correspondingly, the measurements of the natural concentration of carbon dioxide in our atmosphere have shown that the annual increase rate over the past 50 years has been about 100 times faster than previous natural increases, which were due to thermal fluctuations in geological eras. Fig. 1.2 shows the historical fluctuations of the atmospheric carbon dioxide concentration, as reconstructed from ice cores [10]. Global energyerelated CO2 emissions grew by 1.4% in 2017, reaching a historic high of 32.5 gigatonnes (Gt), a resumption of growth after 3 years of global emissions remaining flat [7]; the average concentration of CO2 in the atmosphere also hit a new record high in 2017: 405.0 parts per million (ppm). Fig. 1.3 reports the atmospheric CO2 levels measured at the altitude of 3400 m, at Mauna Loa Observatory, Hawaii, from 2005 to present; the latest data, expressed as dry mole fraction of the number of molecules of carbon dioxide divided by the number of molecules of dry air multiplied by one million ppm, are obtained from direct measurements of the infrared absorption, showed in January 2019 a level of 411 ppm. Despite some debates, there is scientific consensus on two facts, namely that the greenhouse gases are causing a climate change (even if still there is

Figure 1.2 Atmospheric concentration of carbon dioxide during the last three glacial cycles, as reconstructed from ice cores, and current level. Credit: NOAA. Reproduced from NASA, Global Climate Change. Facts. https://climate.nasa.gov/evidence/.

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Figure 1.3 Atmospheric concentration of carbon dioxide from 2005 to present, as obtained by monthly averaged measurements of infrared absorption at the Mauna Loa Observatory, Hawaii. The fluctuations due to the average seasonal cycle were removed from the plot. Credit: NOAA. Reproduced from NASA, Global Climate Change. Carbondioxide. https://climate.nasa.gov/vital-signs/carbon-dioxide/.

no universal agreement about the quantitative magnitude of anthropogenic global warming and on its consequences) and that the CO2 emissions from the combustion of fossil fuels give the largest contribution to the greenhouse gases. Accordingly, policies’ choices made by most governments are aimed at reducing air pollution, ensuring universal energy access, and achieving longterm climate goals, with the central aim to keep a global temperature rise this century well below 2 C above preindustrial levels and to make all efforts to limit the temperature increase even further to 1.5 C. This was the target of the Paris Agreement, entered into force in November 2016 and currently ratified by 185 Parties of 197 Parties to the UNFCCC (United Nations Framework Convention on Climate Change) [11]. The goal of reducing CO2 emissions while still providing the energy services necessary to accommodate the global economic growth may be reached only by reducing substantially the impact of the energy sources, which provide the highest CO2 emissions, namely coal and, more generally, fossil fuels. Electricity is increasing its share among worldwide end uses of energy; industrial electric motor systems, electric mobility, electric heating, and electricity access could lead to a 90% rise in power demand from today to 2040, according to the IEA forecast [9]. It is here that renewable sources of energy can play a substantial role. Fig. 1.4 shows that in 2017, the fraction of electricity generated by renewables reached a record high of 25% (and constituted nearly half of the global additional generation of 380 TWh required to meet the increasing demand); moreover, IEA forecasts that

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Figure 1.4 Electricity generation in 2017, for a total of 25,570 TWh, by different energy sources. PV, photovoltaics. Credit: IEA. Reproduced with modifications from Global Energy & CO2 Status Report 2017. https://www.iea.org/geco/.

the share of generation from renewables could rise from 25% today up to around 40% in 2040 [7]. Today, wind, hydropower, and solar photovoltaics (PV) represent a major source of low-emissions electricity; IEA expects that the increasing competitiveness of solar PV could lead its installed capacity to overcome that of wind before 2025, of hydropower around 2030, and of coal before 2040. Other problems, however, have to be considered; for instance, the wider distribution of PV-generated electricity may create additional requirements for the reliable operation of power systems. Moreover, if electrification certainly reduces local pollution, additional measures may be necessary to control its generation as well as the operation and the decommissioning of the power supply systems from the environmental point of view; otherwise, the risk is that pollution and CO2 emissions simply move upstream from the end-use sectors to power generation, systems’ decommissioning, and recycling. In the present chapter, a brief overview of the status of the research on the fundamental bricks of the solar PV, namely the PV cells, is outlined, without entering into considerations on solar panels and concentrated photovoltaics (CPV).

1.2 Solar photovoltaics: historical notes Solar radiation at the earth’s upper atmosphere totals some 174 PW; the average annual solar radiation is roughly 1.36 kW/m2 and, since about

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30% is reflected back to space, the maximum normal surface irradiance at sea level in a clear day is approximately 1 kW/m2. To be more precise, according to the ASTM G-173 standard [12], which measures solar intensity over the band 280e4000 nm, the total integrated irradiances for the direct and hemispherical tilted (that is, direct beam plus diffuse sky) spectra are 900.1 W/m2 and 1000.4 W/m2, respectively. The latter value is corresponding to the so-called AM1.5 atmospheric condition, namely assuming an absolute air mass 1.5 (solar zenith angle 48.19 ), which is useful to represent the overall yearly average for mid-latitudes. Such a huge amount of power is freely available, and sun’s energy has in fact been used for millennia in various ways, either directly (e.g., by concentrating sun’s rays to make fire and to light sacred fires and torches for religious purposes) or indirectly (e.g., by living in south-facing caves ordat Roman timesdby orienting southwest the largest windows of their bath houses). The sun as an energy source started to be studied in the 18th century, and scientists like Horace de Saussure, John Herschel, and Samuel Pierpont Langley demonstrated that the temperature inside a properly designed glass-covered box could exceed the boiling point of water. Those hot boxes can be considered the prototypes for the solar collectors of recent times. The breakthrough, however, occurred with the discovery of the PV effect by Edmond Becquerel, in 1839 [13], who experimented with metal electrodes immersed in an electrolyte liquid, and with the first “real” solid-state PV cell built in 1883 by Charles Fritts, using junctions formed by coating a plate of copper with selenium and, over it, an ultrathin, nearly transparent, layer of gold [14]. At the beginning of the 20th century, the theoretical aspects of the fundamental photoelectric effect were fully clarified by the 1905 paper of Albert Einstein “On a Heuristic Viewpoint Concerning the Production and Transformation of Light” (he was awarded the 1921 Nobel Prize in Physics for that work) [15] and received experimental proof with the work of Robert Millikan in 1916 [16]. Finally, the 1950s opened the development of PV solar power. The first modern PV cell, based on silicon, was demonstrated by Daryl Chapin, Calvin Fuller, and Gerald Pearson at Bell Laboratories in early 1954 and exhibited a conversion efficiency of 4% (later going up to 11%). Given the high cost and the low efficiency, in the first years the development of PV cells was focused on the satellites’ applications; Vanguard 1, launched in 1958 in the frame of a program managed by the United States Naval Research Laboratory, was the first satellite to use radios powered by solar

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energy (less than 1 W). Few years later, the first commercial communications satellite Telstar 1, designed and built by a team at Bell Telephone Laboratories and launched by NASA in 1962, was powered by 3600 solar cells, for a total 14 W. Starting in the 1970s, with the work of Elliot Berman at Exxon, the improvements in materials and processes made the price of silicon solar cell continuously decreasing, down to a fraction of US$ per generated electrical watt after 2010. It made the use of cells in large solar panels affordable and practical for both residential and nonresidential applications. The decline of price of silicon PV cells from 1977 to 2015 is shown in Fig. 1.5 [17]. In the last decade, PV research and development has continued, focusing onto new materials and advanced cell designs. In particular, researchers have been looking for materials that can perform at similar or better levels than silicon (in its various types, namely, monocrystalline, polycrystalline, and amorphous silicon) but can be easily available and cheaper to produce. These new materials include organics, IIIeV semiconductors, thin-film semiconductors (e.g., CdTe and copper indium gallium diselenide (CIGS)),

Figure 1.5 Price history chart of crystalline silicon solar cells in US$ per watt from 1977 to 2015. PV, photovoltaics. Reproduced from Wikipedia Commons, Price History of Silicon PV Cells Since 1977, 2015. https://commons.wikimedia.org/wiki/File:Price_history_of_ silicon_PV_cells_since_1977.svg.

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perovskites, quantum dots, carbon nanotubes; their basic principles are quickly summarized in the following section.

1.3 Main types of solar cells The evolution of solar cells’ technologies, briefly introduced in the previous section, is usually divided into three generations. The first generation is mainly based on monocrystalline or polycrystalline silicon wafers. This generation is well established now and is commercially mature, covering about 80% of the solar market. The efficiency of these cells is about 22%e25% in the laboratory (Fig. 1.6), mainly limited by the spectral response of silicon single-junction semiconductor. The second generation is based on thin-film semiconductors, mainly amorphous silicon, cadmium telluride (CdTe), and CIGS. The significant material amount reduction and the different production technologies (like vacuum deposition processes) make these cells much cheaper than silicon wafer technologies, with new features like flexibility, which allow for peculiar applications, although their efficiency is slightly lower than the one of traditional cells. Finally, the third generation of solar cells is based on the possibility to exploit novel materials and architectures, aiming at very cheap productions and/or very high efficiencies. Among them, multijunction solar cells (MJSCs), dyesensitized solar cells (DSSCs), quantum dotesensitized solar cells, organic

Figure 1.6 Chart of the best efficiencies achieved with research photovoltaic (PV) cells, plotted from 1976 to 2018. Credit: National Research Energy Laboratory. https://www. nrel.gov/pv/cell-efficiency.html.

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solar cells (OSCs), and perovskite solar cells (PSCs) are some examples. Each of them has specific advantages and limits, regarding not only cost, stability, and efficiency but also peculiar properties like flexibility, transparency, etc. The highest efficiency reported to date in the laboratory is as high as 47% in MJSC. An almost exhaustive list of the many types of solar cells developed in the laboratories, both in academies and industries, is presented in Fig. 1.6, which shows the time evolution of the best efficiencies achieved worldwide. This plot is the 2018 version (issued in December 2018) of the chart regularly maintained by NREL, USA, with the highest conversion efficiencies for research cells from 1976 to the present [18]. Devices included in this chart have efficiencies that are confirmed by independent, recognized test labs and are reported on a standardized basis; the most recent world record for each technology is highlighted along the right edge in a flag that contains the efficiency and the symbol of the technology. The company or group that fabricated the device for each most recent record is bolded on the plot. The plot in Fig. 1.6 lists five families of PV cells (i.e., multijunction cells, single-junction gallium arsenide cells, crystalline silicon cells, thinfilm technologies, emerging PV), subdivided in 28 subcategories. A quick overview of these structures is given in the following subsections.

1.3.1 Multijunction solar cells MJSCs are currently the most efficient way to convert sunlight into electricity (Fig. 1.7). The basic idea is to combine different semiconductor

Figure 1.7 (A) The structure of a multijunction solar cell. (B) Graph of spectral irradiance versus wavelength over the AM1.5 solar spectrum, together with the maximum electricity conversion efficiency for every junction as a function of the wavelength. By NcouniotdFraunhofer Institute for Solar Energy Systems, CC BY-SA 3.0.

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materials, with different band gaps. Since photons with energy below the band gap are transmitted without losses, the architecture of such a device has the highest band gap junction first and the lowest band gap material last. As the photons have to pass through the cell to reach the proper layer to be absorbed, transparent conductors need to be used to collect the electrons being generated at each layer. In the realization of a monolithical cell, the layers are mechanically and electrically connected. Therefore, the choice of the materials for each layer must fit specific requirements, not only based on a suitable repartition of the solar light spectral interval but also related to their mechanical and electrical properties. To obtain the best device performance, lattice matching and current matching are desired. Typically, these cells are grown layer-by-layer and for optimal growth and resulting crystal quality, the crystal lattice constant should be closely matched, resulting in lattice-matched devices. A great degree of mismatch or other growth imperfections can lead to crystal defects causing degradation in electronic properties. Since each subcell is connected electrically in series, the same current flows through each junction. Therefore, suitable band gaps must be chosen such that the design spectrum will balance the current generation in each of the subcells, achieving current matching. In Fig. 1.7B, the AM1.5 solar spectral irradiance is plotted as a function of wavelength, together with the maximum conversion efficiency for every junction, which is directly related to the number of photons available for conversion into photocurrent. Finally, strong absorption coefficients, high minority carrier lifetimes, and high mobilities are desired, justifying the choice of materials typically used for MJSCs: InGaP for the top subcell (Eg ¼ 1.8e1.9 eV), InGaAs for the middle subcell (Eg ¼ 1.4 eV), and germanium for the bottom subcell (Eg ¼ 0.67 eV). In 2014, an MJSC established the world record efficiency at 46% for the direct conversion of sunlight into electricity and was developed by Soitec and CEA-Leti, France, together with the Fraunhofer Institute for Solar Energy Systems ISE, Germany: it was a four-junction cell. Each of its subcells converts exactly one quarter of the incoming photons in the wavelength range between 300 and 1750 nm into electricity. The record was obtained in concentrated PV conditions, by using a Fresnel lens, at a concentration of 508 suns. The perfect current matching condition obtained, thanks to the exact distribution of the photons among the four subcells, was achieved by a precise tuning of the composition and thicknesses of each layer inside the cell structure [19].

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1.3.2 Gallium arsenide single-junction solar cells Gallium arsenide (GaAs) solar cells are considered as a separate family of PV devices, although they are made as thin-film layers deposited on a supporting substrate. GaAs is one of the most commonly used IIIeV semiconductor materials. The high electron mobility, direct band gap, high quality of the epitaxial growth process, and excellent surface passivation by AlGaAs allowed GaAs to reach the world record efficiency for single-junction thin-film devices in December 2018, up to 29.1% (Fig. 1.6), significantly higher than silicon (record held by Alta Devices). The possibility to realize efficient thin-film solar cells allows the overall weight of the solar panel to be kept low, with the additional property of being flexible. Noteworthy, GaAs is naturally resistant to damage from moisture, radiation, and ultraviolet light. These properties, together with lightweight, make GaAs an excellent choice for aerospace applications where increased UV and exposure to radiation are an issue. Furthermore, most solar cells are inefficient in weak illumination conditions (low light intensity), but not high-quality GaAs. This feature, together with the property of a wider band gap with respect to silicon, makes GaAs very suitable also for indoor applications under LEDs and fluorescent light excitation. In Chapter 10, we will see an overview of solar cells for indoor applications [20], with a focus on some novel alternative approaches, based on organics and perovskites.

1.3.3 Silicon solar cells The use of silicon in PV technologies has been already introduced in previous paragraphs as the first generation of solar cells, and it will be discussed in depth in Chapter 2 of this book [21]. Silicon PV is considered as a benchmark: crystalline silicon is the most common material for commercial solar cells, combining affordable costs (Fig. 1.5), good efficiency up to 26% e27% (Fig. 1.6), long-term stability and robustness, together with a solid industrial technology know-how. Silicon has an energy band gap of 1.12 eV, a value that is well matched to the solar spectrum. As we will see in the next section, this value is close to the optimum value for solar-to-electric energy conversion using a single light absorber. Since the band gap is indirect, namely the valence band maximum is not at the same position in momentum space as the conduction band minimum, silicon has a relatively low absorption coefficient. However, with proper light management design, such as the use of surface texturing, antireflection coatings, and rear mirrors,

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efficient light absorption is achieved even with relatively thin wafers (down to w100 mm). On the good side, because of the indirect band gap, radiative recombination is inefficient, which means that the photogenerated electrons and holes in principle can have very long lifetimes. Crystalline silicon solar cells make use of mono- and multicrystalline silicon wafers wire-cut from ingots and cast silicon blocks. An alternative to standard silicon wafer technology is constituted by amorphous or nanocrystalline silicon thin films, which will be described in the next subsection.

1.3.4 Thin-film technologies Thin-film solar cells are considered the second generation and are obtained by depositing one or more thin layers of PV material on a substrate, such as glass, plastic, or metal. Different materials are used in thin-film technologies (Fig. 1.6), such as CdTe, CIGS, and amorphous (a-Si) or micro/nanocrystalline silicon (mc-Si). A common feature of these materials is their direct band gap, which allows a very efficient absorption of light even when their thickness is small. Amorphous silicon has been extensively studied starting with its introduction as a reliable power source of watches, clocks, and calculators in the late 1980s (some examples will be reported in Chapter 10 of this book [20]). On the other hand, CdTe commercial applications have been slightly limited due to the usage of cadmium, which proved to be harmful to both the producer and the consumer. However, CIGS and CdTe are much more promising in terms of energy conversion efficiency than a-Si. A detailed review of thin-film technologies was recently presented by Lee and Ebong [22]. Here we give just a brief summary of the main achievements in the field. Thin-film a-Si material is typically obtained by plasma-enhanced chemical vapor deposition (PECVD) from Si- and H-containing precursor gases, like silane and hydrogen. These amorphous films contain a high amount of hydrogen, as reported for the first time in 1969 by Chittick, Alexander, and Sterling, resulting in a lower defect density and increased photoconductivity. Few years later, in 1975, Le Comber and Spear discovered the possibility of substitutional doping of a-Si:H using phosphine (n-type) or diborane (p-type), which was a fundamental step toward the realization of a solar cell device [23]. Even if a-Si suffers from lower electronic performance compared to crystalline silicon, in particular for solar PV, it can provide some advantages for

Solar cells’ evolution and perspectives: a short review

13

specific and novel applications. For example, amorphous silicon thin films produce savings on silicon material cost. Moreover, thin films are flexible, and PECVD deposition is a low temperature process, which can be done not only on glass but also on plastic substrates, making it a candidate for a roll-to-roll processing technique. Noteworthy, the band gap of thin-film amorphous silicon is about 1.7 eV, corresponding to 730 nm. This puts a limitation to the available region of the solar spectrum that can be absorbed. However, since the whole visible spectral region can be converted in electricity, it makes it a very suitable low-cost PV material for the wide market of indoor PV and for powering Internet of Things (IoT) devices. This emerging and exciting subject of indoor lighting will be widely discussed in Chapter 10 [20]. To shift the spectral absorption band toward the longer wavelengths, micro/nanocrystalline silicon thin films can be realized, which are made of an amorphous silicon matrix containing small silicon crystals. This results in the possibility to combine different silicon thin layers (a-Si/mc-Si) in a multijunction device where the spectral absorption region is controlled by the size of the crystallites in each layer. Following this idea, researchers of the Research Center for Photovoltaics obtained in March 2018 the world record efficiency (14.0%) for a triple junction device based on silicon thin-film technology (Fig. 1.6) [24,25]. The structure of a CIGS solar cell is built on a soda lime glass as the substrate (Fig. 1.8). On top of the glass is the molybdenum contact, then the p-type Cu(InGa)Se2 forms the main junction with n-type CdS, which serves as the buffer layer. An intrinsic zinc oxide layer lies on top of the CdS and, finally, the n-type ZnO:Al layer functions as the front contact. Altogether, the semiconductor used in this solar cell is only 4 mm thick, much lower than in conventional silicon solar cells (about 180 mm). A significant device improvement has occurred since the first thin-film CIGS solar cell was realized by Kazmerski et al. [26] in 1976, having a conversion efficiency of 4.5%. Nowadays, the efficiency record for CIGS thin-film solar cells has been published by researchers at Atsugi Research Center, Solar Frontier K. K., with a value of 22.9% [27]. The main improvement was obtained by modifying the absorber formation, reducing defect density, as suggested by the enhanced photoluminescence performance. As a result, the significant enhancements in open circuit voltage (Voc) and fill factor (FF) led to the achievement of this record-breaking efficiency. However, the same company has announced in January 2019 a new record at 23.35% [28].

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Figure 1.8 (A) Cross-sectional SEM image (left) and schematic structure of a copper indium gallium diselenide (CIGS) solar cell. By L. Kazmerski, National Renewable Energy Laboratory. This work is in the public domain in the United States because it is a work prepared by an officer or employee of the US Government as part of that person’s official duties under the terms of Title 17, Chapter 1, Section 105 of the US Code.

Finally, also the evolution of CdTe solar cell devices followed a significant growth from the first laboratory CdTe cell reported in 1972 by Bonnet and Rabenhorst, who developed a thin-film graded gap CdTe-CdS pen heterojunction solar cell with 6% efficiency. In this field, as for CIGS solar devices, the most advanced industrial equipment and process optimization capabilities are playing the major role, with first solar dominating the past decade, with the highest record efficiency set at 22.1% (Fig. 1.6). Thin-film technologies are an important sector of the PV market. However, some concerns are arising about the shortage and toxicity of materials used in CIGS and CdTe solar cells, both related to material supplying and to their disposal and recycling process. Moreover, the stability issues, in particular for a-Si, are also limiting aspects for market permeation. At the same time, the low-cost manufacturing and continuous rising of efficiencies of crystalline silicon technologies, together with the increasing exploration of alternative approaches in next-generation PV technologies, are seriously affecting the future of these traditional thin-film solar cells. Indeed, their

Solar cells’ evolution and perspectives: a short review

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market share is constantly decreasing after their peak of 17% in 2009, to less than 5% in 2017 [29].

1.3.5 Emerging photovoltaics Besides the traditional approaches previously described, some innovative and emerging technologies have gained attention in the scientific community. A significant part of this book will be focused on these novel materials for PV applications; therefore, we will just mention their main aspects in the following lines. The first real breakthrough in solar cells after silicon was represented by DSSCs, which were first developed by Gr€atzel and O’Regan in 1991 at UC Berkeley [30]. A modern DSSC is composed of a porous layer of titanium dioxide nanoparticles (NPs), covered with a molecular dye that absorbs sunlight, like the chlorophyll in green leaves. The titanium dioxide is immersed in an electrolyte solution, above which is a platinum-based catalyst. Sunlight absorbed by the dye can excite electrons that then flow into the titanium dioxide and then toward the transparent electrode. After flowing through the external circuit, they are reintroduced into the cell on a metal electrode on the back, flowing into the electrolyte, which transports the electrons back to the dye molecules [31]. A detailed description of the evolution and latest advances in DSSC is given in Chapter 4 of this book [32]. DSSCs, however, are still at their development stage. Some novel advances include codoping of TiO2 layer with quantum dots [33], plasmonic metal NPs [34], or carbon-related materials (nanotubes, graphene) [35] to increase the current transport properties, using solid-state electrolytes for better temperature response and changing the doping of the TiO2 to better match it with the electrolyte being used. Another significant emerging field is that of PSCs. PSCs are based on an absorber with ABX3 perovskite crystal structure where “A” and “B” are cation elements and “X” is an anion element. Originally, the terminology of “perovskite” materials was accepted in the first half of the 19th century, when the group of Lev Perovski and Gustav Rose discovered a minerald CaTiO3 with ABX3 structure. Hybrid organiceinorganic halide PSCs have risen to stardom owing to the peculiar characteristics of the halide perovskite absorber, such as high charge carrier mobility, broad and strong optical absorption, long free carrier diffusion length, low exciton binding energy, as well as their cost-effective and easy solution process manufacture. Noteworthy, PSCs have achieved remarkable progress with power

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conversion efficiency (PCE), developing from 3.8% in 2009 to 23.7% in 2018. The main aspects and perspectives of this technology will be presented in Chapter 5 of this book [36]. Organic or polymer solar cells (OSCs) have also attracted tremendous attention owing to their great potential for large-area production with low-cost materials and simple solution processing method. OSCs usually consist of an electron- or hole-blocking layer on top of an indium tin oxide (ITO) conductive glass, followed by an electron donor and an electron acceptor (in the case of bulk heterojunction solar cells), a hole or electron blocking layer, and a metal electrode on top. Among them, ternary OSCs featuring multiple light harvesting materials in one active layer have emerged as one of the most promising strategies to improve PV performance. The PCE of ternary OSCs has approached 14.62%. In Chapter 3 of this book, the development of ternary OSCs based on different systems is summarized, as well as the main working mechanisms and the potential research directions and perspectives [37]. Finally, all-oxide solar cells promise to provide vacuum-free and low-cost solutions, based on Earth’s abundant and environmentally friendly materials. While oxides can exhibit high extinction coefficient in the visible and near-infrared region (NIR) spectral regions, guaranteeing full absorption of sunlight, the main factor limiting efficiency in such kind of pen junction devices is the low hole mobility in the p-type oxide, which represents the main challenge to be overcome to make this technology competitive. Chapter 6 illustrates the latest results in the field, including integration of nanowire geometries as viable solution toward fast charge transport and collection [38].

1.4 Efficiency of solar cells The basic model of a traditional PV cell is represented by a semiconductor pen junction (Fig. 1.9A), with metallic contacts deposited on the top and bottom [39]. A detailed description of this device is reported in Chapter 7 of this book [40]. In summary, due to the diffusion of majority carriers, a depletion region is established across the pen junction, with the building up of an electric field and potential difference between the two regions. Photons incoming into the cell’s active area, corresponding to this depletion region, have the potential to give their energy (hn) to electrons and to generate if hn  Eg (band gap energy of the semiconductor), an electrone hole pair. This pair is suddenly divided by the present electric field, resulting

Solar cells’ evolution and perspectives: a short review

(A)

(B)

sunlight

p+

17 l

Depletion Layer

Electron flow

photonic electron excitation photonic electron excitation

n+

Hole flow Electric current flow

Ec Ef Ev

dark

V

illuminated

Rload

Figure 1.9 (A) Illustration of a solar cell device structure in the form of pen diode with external load. (B) JeV characteristic of the solar cell in the dark and under illumination. Reproduced from P.S. Priambodo, N.R. Poespawati, D. Hartanto, Solar Cell, in: Kosyachenko (Ed.), Sol. Cells - Silicon Wafer-Based Technol., IntechOpen, 2011.

in a generated photocurrent (IL). In practice, the JeV characteristic of a solar cell under illumination (Fig. 1.9B) is obtained by shifting downward by IL the JeV characteristic of the same diode in the dark. Four important parameters can be defined when dealing with solar cells. The first parameter is the short circuit current (ISC), which is measured when the output terminal is shorted, or V is equal to 0. This value of output current ISC represents the maximum current provided by the solar cell at a certain illumination level. The second parameter is the open circuit voltage (VOC), which is measured when the output terminal is opened, or I is equal to 0. This value of output voltage VOC represents the maximum voltage of solar cell at a certain illumination level. In Fig. 1.9B, the maximum delivery output power (PMP) is represented by the area of product VMP by IMP as the maximum possible area included within the curve and the axes. The third parameter is the fill factor (FF) that represents the ratio between the maximum power PMP and the product VOC times ISC: PMP VMP $IMP ¼ (1.1) VOC $ISC VOC $ISC This parameter gives an idea of how squared the output characteristic of the cell is, and for real devices it is typically between 0.7 and 0.85. Finally, the fourth parameter is the efficiency (h). The efficiency of a solar cell is defined as the ratio of the electrical energy output to the total incoming sunlight energy striking the surface of the cell. Many parameters affect the determination of the overall efficiency, which depends on the architecture and materials of the cell, their sensitivity to the different wavelengths of the solar FF ¼

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spectrum, the intensity of the incident light, the working temperature, the surface reflectivity and texturing, the efficiency in the generation of electronehole pairs, and the capability of subsequent extraction of these charge carriers without losses. The efficiency can be written as VMP $IMP VOC $ISC $FF ¼ (1.2) Pin Pin The four key parameters described above are controlled by the materials that constitute the solar cell, and some discussions are presented in other chapters of this book (e.g., see organic SC in Chapter 3 [37], dyesensitized SC in Chapter 4 [32], and perovskite SC in Chapter 5 [36]). Furthermore, the parameters also depend on the spectral properties and intensity of light illumination (e.g., see the discussion on indoor lighting conditions in Chapter 10 [20]). A detailed analysis of the factors affecting energy conversion efficiency was presented in a landmark paper by William Shockley and Hans Queisser in 1961 [41]. Among the factors concerning energy losses, firstly we can consider the losses due to blackbody radiation. In the case of solar cells at standard temperature and pressure, this loss accounts for about 7% of the power. The second factor is due to possible electronehole recombination events. In silicon, for example, this accounts for another 10% of the power. However, the third and dominant energy loss mechanism is related to the spectral response of the solar cell active material. This is because the photons with energy (hn) below the energy gap of the semiconductor are not able to generate electronehole pairs and are lost. Conversely, photons with more energy than the band gap initially eject an electron to a state high above the band gap, but this extra energy is lost through collisions in internal relaxation paths. This lost energy turns into heat in the cell, which has the side effect of further increasing blackbody losses [42]. Combining all these factors, the maximum theoretical efficiency for a single-band gap material, like conventional silicon cells, is about 34%. An analysis of the AM1.5 spectrum shows that the best balance is reached at about 1.1 eV (1100 nm, in the near-infrared), which happens to be very close to the natural band gap in silicon (Fig. 1.10) [43]. To overcome the ShockleyeQueisser limit, a significant efficiency increase can be obtained by employing different semiconductor materials to cover different regions of the solar spectrum. These MJSCs, as described in the previous section, combined with concentrated solar light geometry, h¼

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Figure 1.10 The ShockleyeQueisser limit for the efficiency of a single-junction solar cell under unconcentrated sunlight at 273 K. This calculated curve uses actual solar spectrum data, and therefore the curve is wiggly from IR absorption bands in the atmosphere. This efficiency limit of w34% can be exceeded by multijunction solar cells. Reproduced from Wikipedia Commons, Shockley Queisser Full Curve. https://commons. wikimedia.org/wiki/File:ShockleyQueisserFullCurve.svg.

established efficiency records above 46%, at expenses of complexity and cost. MJSCs may find applications in special sectors like satellites and other space devices, where size, weight, and other parameters are fundamental factors to be considered, and the cost of the solar panels is not an issue. On the other hand, commercial everyday applications and the development of new markets and needs, like flexible, wearable, short-lasting devices, pushed the research to the development of solar cells where the low cost or some special functionalities are the driving features while accepting a lower efficiency. At the same time, however, new innovative approaches and strategies, mainly based on optical concepts, are being developed for increasing the efficiency of solar cell devices. Among them, the most important are light trapping strategies, plasmonics, and spectral conversion, which will be shortly described in the following.

1.4.1 Light trapping strategies For solar cells, a significant contribution to energy loss is the reflection at the front surface. Therefore, antireflective (AR) coatings are routinely used, but better results can be achieved by proper texturing the front surface: it is thus possible to additionally enhance the omnidirectionality, i.e., the capacity of trapping the photons from various incident directions through the scattering effect, induced by the micro- or nanostructures fabricated on the surface

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[44e49]. This is particularly important for thin-film solar cells, where the short optical path is not enough to absorb all the incoming solar radiation, and it can be useful also for other emerging solar technologies. Conventional light trapping approaches are based on textured rear/front surfaces, to diffuse light and increase the optical path length within the semiconductor absorber layer [50e52]. Alternative strategies have been also investigated, employing nano-/microstructures with dimensions comparable or smaller than the illuminating wavelengths, such as diffraction gratings [53], microlenses [54e56], Mie resonators [57,58], and plasmonic metal NPs [59e61]. Nonetheless, the main drawback of texturing the semiconductor surface is the fact that it increases roughness (hence surface area) and defect density in the PV material, which deteriorates the cells’ electrical transport via the increase of charge carrier trapping and recombination. A detailed description and analysis of the advantages and challenges related to different light trapping strategies is presented in Chapter 9 of this book [41,62].

1.4.2 Plasmonics As previously observed, surface texturing in solar cells may increase loss mechanisms related to surface recombination and charge carrier trapping at defect states. An alternative and more efficient approach to obtain efficient light concentration and trapping is based on plasmonic structures, mainly constituted by patterned metal films or NPs. Plasmonics has demonstrated to be very efficient for redirecting and enhancing the incident light beam, especially in the field of PV [63e65]. A comprehensive updated summary of the most important results and perspectives in this exciting field is reported in our recent review paper [66]. As shown in Fig. 1.11, plasmonic nanostructures have been used at least in three different configurations: on the surface of the cell (a) [67e71], embedded into the semiconductor active layer (b) [72e73], or at the bottom interface between the semiconductor and the metal contact (c), in the form of scattering NPs [60,74,75] or surface plasmonic gratings [76e79]. The increase of the cell efficiency is due to different physical principles. In the (a) and (c1) configurations, the multiple and high-angle scattering of light causes an increase of the effective optical path length inside the cell. In the (b) configuration, a strong localization and enhancement of the nearfield in the proximity of the particles causes the creation of additional electronehole pairs within the semiconductor. Finally, in the (c2)

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Figure 1.11 Plasmonic light trapping configurations for thin-film solar cells. (A) Metal nanoparticles on the surface of the solar cell. (B) Metal nanoparticles embedded in the semiconductor. (C1 and C2) Metal nanoparticles or periodic arrays at the metal/ semiconductor interface. Reproduced from F. Enrichi, A. Quandt, G.C. Righini, Plasmonic enhanced solar cells: summary of possible strategies and recent results, Renew. Sustain. Energy Rev. 82 (2018) 2433e2439. https://doi.org/10.1016/j.rser.2017.08.094.

configuration, a metal grating at the back surface can couple light to surface plasmon polaritons, which propagate in the plane of the semiconductor layer. As a general rule, when scattering is the main occurring process, the maximum efficiency is achieved typically by using 100e200 nm metal NPs (the optimum size depends on the metal and on the optical properties of the surrounding material). Instead, when localized near-field is involved (as in the case of particles embedded into the active area of the semiconductor), the best results are obtained by using 10e20 nm NPs. Larger NPs induce detrimental absorption losses and heating. Noteworthy, the spectral response of the metallic NPs is also a fundamental aspect that must be kept in consideration. Indeed, since absorption losses are significant in the resonant spectral region, a filtering effect occurs when the NPs are deposited on the surface of solar cells, and the increase in efficiency may be possible only for wavelengths longer than the resonance range and for active materials not absorbing in that resonance region. Typical metallic NPs give the best increment in the red spectral region (600e700 nm), while the blue-green regions are penalized. However, the application of the NPs at the back of the solar cells can avoid this filtering effect, and it has the further advantage of providing a broadband spectral enhancement. Recently, it was demonstrated that high-permittivity dielectric NPs can mimic some peculiar properties of plasmonic metallic particles, opening the route to light manipulation at the nanoscale by using dielectric low-loss materials [55,80e83]. This line of research is expected to bring significant impact also to the field of solar cells in the near future, and the combination

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of hybrid metaledielectric nanostructures has been proposed as the best candidate for achieving the highest efficiency values, surpassing all-metal and all-dielectric scatterers [84]. Specific examples and strategies on light trapping by plasmonic NPs are discussed extensively in Chapter 8 of this book [85].

1.4.3 Spectral conversion As already discussed, the major limitation to the efficiency of singlejunction solar cells is intrinsically related to the limited spectral response of such a device. In fact, the system is optimized only for photon energies (hn) close to the band gap (Eg) of the semiconductor material. Higherenergy photons provide an energy excess, which is lost by internal relaxation mechanisms, while lower-energy photons pass through the material without being absorbed. Currently, the best strategy to increase the electrical conversion efficiency of the solar radiation is to use different band gap semiconductor materials in an MJSC, as previously described. However, an alternative and cheaper strategy is provided by applying passive spectral conversion layers on the surface or on the back-face of the solar cell for down-converting higher-energy photons and up-converting lowerenergy photons (Fig. 1.9) [86,87]. An important aspect of this approach is that spectral converters are readily applicable to existing solar cells, with only few changes and almost without affecting the production process. Moreover, the spectral converters and the solar cells can be optimized independently. As depicted in Fig. 1.12, the spectral conversion approach could, in principle, significantly improve the total available energy for electrical generation. According to theoretical calculations, the maximum efficiency which could be obtained by concentrating the whole solar spectrum into a single narrow emission at the semiconductor band gap would be higher than 80% [88]. A detailed analysis of the down-conversion and upconversion potential was done by Trupke et al. in two parallel papers. The maximum efficiency for a solar cell covered with an ideal downconversion layer could be as high as 38.6% [89], while the application of an ideal up-converter on the backside of a solar cell could result in a maximum efficiency of 47.6% [90]. Both these results were calculated by using nonconcentrated sunlight; they clearly show that the development of optimized spectral conversion layers would be a cheap high-gain solution for PV solar cell technologies.

Solar cells’ evolution and perspectives: a short review

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Figure 1.12 AM1.5 spectrum showing the fraction that is currently absorbed by a thick silicon device and the additional regions of the spectrum that can contribute toward up- and down-conversion. Reproduced with permission from B.S. Richards, Enhancing the performance of silicon solar cells via the application of passive luminescence conversion layers, Sol. Energy Mater. Sol. Cells 90 (2006) 2329e2337. https://doi.org/10.1016/j. solmat.2006.03.035.

Three luminescence processes, defined as quantum cutting, downshifting, and up-conversion, which are sketched in Fig. 1.13, are currently explored for reducing the spectral mismatch [91]. In a typical downshifting process, upon excitation with a high-energy photon, a nonradiative relaxation takes place followed by the radiative emission of a lower-energy photon. This process does not increase the total number of available

Down-shifting

Down-conversion Visible Up-conversion

Photon Energy

UV

NIR

Figure 1.13 Mechanisms of photon luminescent conversion processes: down shifting, down-conversion, and up-conversion [91]. Only two-photon processes are shown for illustrative purposes.

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photons, but it brings them from the UV region, where they cannot be exploited by the cell, to the region of maximum efficiency, close to the semiconductor band gap. Among materials with excellent downshifting properties, semiconductor quantum dots are very suitable for increasing the efficiency of solar cells. In fact, they have a very broadband absorption range, with maximum in the UV region, high quantum efficiencies, and narrow emissions, which can be finely tuned by proper control of their size and shape. An example of this application is discussed in Chapter 13 of this book [92]. An alternative process for better exploiting high-energy photons is the quantum cutting process, which splits the energy of an absorbed photon by emitting two (or more) lower-energy photons. Trivalent rare earth ions are the prime candidates to achieve efficient spectral conversion because of their rich energy-level structure (Dieke diagram) and long excited state lifetimes that allow for facile photon management. Noteworthy, our group paid significant attention to the development of spectral conversion layers for solar cells based on glass ceramic materials, as also reported in Chapter 11 of this book [93]. A comprehensive review of Yb3þ and RE3þ (RE ¼ Ce, Pr, Nd, Eu, Eu, Tb, Dy, Ho, Er, Tm) codoped materials for down-conversion is presented in Chapter 12 of this book [94]. The optical emission of Yb3þ is particularly interesting for solar cells because it occurs in the NIR at 980 nm, which fits well with the band gap of crystalline silicon. Both downshifting and quantum cutting processes are able to make effective use of the high-energy portion of the solar spectrum but with different quantum efficiencies. The theoretical internal quantum efficiency for downshifting is always less than 100%, whereas it exceeds 100% and approaches 200% for quantum cutting [95e97]. It has, however, to be remarked that the near-infrared emission intensity does not increase at the same rate and that the absolute quantum efficiency (i.e., the ratio of emitted photons to absorbed photons, measurable in an integrating sphere detection system) is much lower [98]. Finally, the up-conversion process can convert two (or more) incident low-energy photons into a single higher-energy photon. Some theoretical modeling about this process is presented in Chapter 7 of this book [40]. Note that the downshifting and quantum cutting layers are generally placed on the front surface of the solar cell, allowing the down-converted photons to be absorbed by the solar cell. Instead, the up-conversion material is

Solar cells’ evolution and perspectives: a short review

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typically placed on the rear face of the solar cell, to harvest the subband gap spectrum of sunlight.

1.5 Challenges and future prospects Achieving a climate-neutral economy is currently one of the longterm targets of most of the countries worldwide. The needed strategy includes, among others, two key actions: energy efficiency and deployment of renewables. Wind and solar PV are definitely a major source of lowemissions electricity, which can help to respond to the expected growing power demand for electric mobility, electric heating, and electricity access from today to 2040. Thanks to the continuous advances in the solar cells’ materials and technologies, and the consequent development of efficient and cheap solar panels, the competitiveness of solar PV is expected to push the PV installed capacity beyond that of wind before 2025, past hydropower around 2030 and past coal before 2040 [9]. Here, only the status of the research on discrete solar cells has been described; the solar industry, however, is mostly based on solar panels and sun-concentrating structures; the interested reader is referred to the 2017 report by Fraunhofer ISE and NREL on the status of CPV technology [99]. Remaining at the device level, the search for more efficient materials and cell structures keeps going; for instance, among materials, even if silicon appears to be capable to remain the market dominator for many years ahead [21], perovskites are receiving a great deal of attention as a potentially cheaper and more efficient alternative [36]. A record efficiency of 18% on a 1.2 cm2 and of 12.1% for a large 16 cm2 single PSC was recently achieved by Ho-Baillie and colleagues at UNSW in Australia [100]. A world record of 23.35% for CIGS thin-film solar cell efficiency was announced on January 2019 by the Japanese Solar Frontier company [28]; it is worth to note that CIGS modules have the additional advantage of requiring up to 60% less energy to produce than crystalline silicon modules. OSCs have also attracted tremendous attention owing to their great potential for largearea production with low-cost materials and simple solution processing method [37], and so did DSSCs, which can be manufactured using rollto-roll processing and are prone to easy integration into building architectures [32]. Staying with silicon, several attempts are being pursued to increase the efficiency of a cell, e.g., through a greater light trapping capability or a better exploitation of the solar spectrum [20,40,62,85,92e94].

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The combination of two materials in the so-called tandem structure may offer an additional degree of freedom in the design of a PV device: organice inorganic or inorganiceinorganic structures provide new paths to the achievement of high efficiency without affecting too much the device cost [101e103]. It is known that significantly higher conversion efficiencies can be reached with multijunction cells based on semiconductors like gallium and indium, but their cost is high due to the raw materials’ cost and manufacturing complexity. The use of nanowires may allow to significantly reduce material needs without compromising absorption or performance [104]. Another possibility to reduce the costs of PV IIIeV multijunction systems is to leverage the efficiency and concentration: NREL demonstrated w46% efficiency with a four-junction IMM (inverted metamorphic multijunction) solar cell, which uses a compositionally graded buffer to incorporate nearly perfect single-crystal layers (GaInPeGaInAseGaInAs) with different crystal lattice parameters [19]. While the use of solar cells in home appliances, personal devices, and IoT does not introduce any special problem, the growing deployment of solar PV electricity-generating systems may create additional requirements for the reliable operation of power systems. On one side, a problem of energy storage exists, given the intermittent production of electricity (abundant when it is sunny, reduced when it is cloudy, none at all when it is dark); it implies the search for novel storage materials, especially novel battery materials. On another side, domestic renewable production, which is going to increase continuously, may revolutionize the current distribution systems and regulatory frameworks that were primarily designed for a centrally controlled reality. Affordability, reliability, and sustainability, which must be guaranteed in the future scenario, are closely interlinked: it is necessary to find a trade-off among various issues, such as the efficient integration of the different electric supplies into the existing grids, the balance with the energy demand, the fragmentation of the solar renewable industry, and the need of flexible but effective regulatory controls. Finally, the target of sustainability must encompass environmental, social, and economic aspects [1]; for instance, as mentioned in a report of European Commission, the growth in renewables is not without challenges in other areas. It may affect in particular the agricultural sector, as competition for land use may put pressure on food prices, but may also put stress on cities, which were built around fossil fuels and now may need to revise infrastructure and urban planning systems [105].

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In this complex frame, however, R&D on solar cells’ physics and technology remains a main actor, and its future is definitely sunny.

Acknowledgments The authors acknowledge the support of the Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi. F.E. acknowledges the support of VINNOVA (Swedish Innovation Agency), under a Vinnmer Marie Curie IncomingdMobility for Growth Programme (project “Nano2solar” Ref. N. 2016-02011).

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[91] A. Shalav, B.S. Richards, M.A. Green, Luminescent layers for enhanced silicon solar cell performance: up-conversion, Sol. Energy Mater. Sol. Cells 91 (2007) 829e842, https://doi.org/10.1016/j.solmat.2007.02.007.  [92] A. Flores-Pacheco, M.E. Alvarez-Ramos, A. Ay on, Utilization of quantum dots on silicon solar cells, in: F. Enrichi, G.C. Righini (Eds.), Sol. Cells Light Manag. Mater. Strateg. Sustain., Elsevier, 2019. [93] A. Chiappini, L. Zur, F. Enrichi, B. Boulard, A. Lukowiak, G.C. Righini, M. Ferrari, Glass ceramic for frequency conversion, in: F. Enrichi, G.C. Righini (Eds.), Sol. Cells Light Manag. Mater. Strateg. Sustain., Elsevier, 2019. [94] Y. Katayama, S. Tanabe, Downconversion for 1mm luminescence in lanthanide and Yb3þ-codoped phosphors, in: F. Enrichi, G.C. Righini (Eds.), Sol. Cells Light Manag. Mater. Strateg. Sustain., Elsevier, 2019. [95] A. Bouajaj, S. Belmokhtar, M.R. Britel, C. Armellini, B. Boulard, F. Belluomo, A. Di Stefano, S. Polizzi, A. Lukowiak, M. Ferrari, F. Enrichi, Tb3þ/Yb3þ codoped silicahafnia glass and glass-ceramic waveguides to improve the efficiency of photovoltaic solar cells, Opt. Mater. 52 (2016), https://doi.org/10.1016/j.optmat.2015.12.013. [96] F. Enrichi, C. Armellini, S. Belmokhtar, A. Bouajaj, A. Chiappini, M. Ferrari, A. Quandt, G.C. Righini, A. Vomiero, L. Zur, Visible to NIR downconversion process in Tb3þ-Yb3þ-codoped silica-hafnia glass and glass-ceramic sol-gel waveguides for solar cells, J. Lumin. 193 (2018) 44e50, https://doi.org/10.1016/ j.jlumin.2017.08.027. [97] L. Zur, C. Armellini, S. Belmokhtar, A. Bouajaj, E. Cattaruzza, A. Chiappini, F. Coccetti, M. Ferrari, F. Gonella, G.C. Righini, E. Trave, A. Vomiero, F. Enrichi, Comparison between glass and glass-ceramic silica-hafnia matrices on the down-conversion efficiency of Tb3þ/Yb3þ rare earth ions, Opt. Mater. 87 (2019) 102e106, https://doi.org/10.1016/j.optmat.2018.05.008. [98] D. Qianqian, F. Qin, H. Zhao, Z. Zhang, W. Cao, Absolute quantum cutting efficiency of Tb3þ-Yb3þ co-doped glass, J. Appl. Phys. 114 (2013) 213513. https://doi. org/10.1063/1.4838435. [99] I.S.E. Fraunhofer, NRL, Current Status of Concentrator Photovoltaic (CPV) Technology, 2017. https://www.ise.fraunhofer.de/content/dam/ise/de/documents/ publications/studies/cpv-report-ise-nrel.pdf. [100] D. Bradley, Record-breaking Perovskite Solar, Materials Today, December 22, 2016. https://www.materialstoday.com/energy/news/recordbreaking-perovskite-solar/. [101] X. Jia, B. Puthen-Veettil, H. Xia, T.C.J. Yang, Z. Lin, T. Zhang, L. Wu, K. Nomoto, G. Conibeer, I. Perez-Wurfl, All-silicon tandem solar cells: practical limits for energy conversion and possible routes for improvement, J. Appl. Phys. 119 (2016) 233102, https://doi.org/10.1063/1.4954003. [102] A.J. Blanker, P. Berendsen, N. Phung, Z.E.A.P. Vroon, M. Zeman, A.H.M. Smets, Advanced light management techniques for two-terminal hybrid tandem solar cells, Sol. Energy Mater. Sol. Cells 181 (2018) 77e82, https://doi.org/10.1016/ j.solmat.2018.02.017. [103] J. Park, V.A. Dao, S. Kim, D.P. Pham, S. Kim, A.H.T. Le, J. Kang, J. Yi, High efficiency inorganic/inorganic amorphous silicon/heterojunction silicon tandem solar cells, Sci. Rep. 8 (2018) 15386, https://doi.org/10.1038/s41598-018-33734-y. [104] NanoTandem, Nanowire-Based Tandem Solar Cells. https://nano-tandem.ftf.lth.se/ index.php. [105] European Political Strategy Centre, 10 Trends Reshaping Climate and Energy, 2018. https://ec.europa.eu/epsc/sites/epsc/files/epsc_-_10_trends_transforming_climate_ and_energy.pdf.

CHAPTER TWO

Silicon solar cells: materials, technologies, architectures Lucia V. Mercaldo, Paola Delli Veneri Italian National Agency for New Technologies, Energy and Sustainable Economic Development (ENEA), Portici Research Centre, Portici, NA, Italy

2.1 The photoactive materials A solar cell in its most fundamental form consists of a semiconductor light absorber with a specific energy band gap plus electron- and holeselective contacts for charge carrier separation and extraction. Silicon solar cells have the advantage of using a photoactive absorber material that is abundant, stable, nontoxic, and well understood. In addition, the technologies, both the crystalline silicon (c-Si) and the thin-film Si-based, can rely on solid know-how and manufacture equipment, having benefited also from the microelectronics industry sector along its historical development.

2.1.1 Crystalline silicon The light absorber in c-Si solar cells is a thin slice of silicon in crystalline form (silicon wafer). Silicon has an energy band gap of 1.12 eV, a value that is well matched to the solar spectrum, close to the optimum value for solar-to-electric energy conversion using a single light absorber. Its band gap is indirect, namely the valence band maximum is not at the same position in momentum space as the conduction band minimum. As a consequence, silicon has a relatively low absorption coefficient near the band edge (the nearinfrared region). However, with proper light management design, such as the use of surface texturing, antireflection coatings, and rear mirrors, efficient light absorption is achieved even with relatively thin wafers (down to w100 mm, while the typical thickness in c-Si solar cell production lines is currently 180 mm [1]). On the good side, because of the indirect band gap, radiative recombination is inefficient, which means that the photogenerated electrons and holes in principle can have very long lifetimes. The dominant intrinsic recombination mechanism in silicon is thus the Auger Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00002-1

© 2020 Elsevier Ltd. All rights reserved.

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contribution, in which three charge carriers are involved, one electron that recombines with a hole by transferring the energy difference to a second free electron or a hole (this energy is subsequently lost as heat). In practice, carrier recombination in silicon can be affected by crystallographic defects or impurities. However, in recent years, the electronic quality of crystalline silicon wafers has improved to such a point that solar cell advances now rely essentially on improvements in surface passivation and in the contacting schemes. Passivation is applied to minimize the recombination of photogenerated carriers at the silicon surface. This is achieved by chemical means, i.e., by providing a termination (usually hydrogen) to the detrimental surface dangling bonds (broken SieSi bonds), and by a field effect mechanism, namely providing a built-in electric field that repels one charge carrier type from the interface (because electrons and holes are needed together for recombination to take place). Effective surface passivation is usually achieved by dielectric materials such as silicon oxide, silicon nitride, silicon carbide, aluminum oxide and by hydrogenated amorphous silicon [2]. On the practical side, c-Si solar cells make use of mono- and multicrystalline silicon (mc-Si) wafers, wire-cut from ingots and cast silicon blocks, respectively. It is estimated that mc-Si wafers have a market share of 52% in the silicon solar cell manufacturing industry today, coming from a 60% versus 40% for mono-Si in 2017 [1]. The most common method for the growth of single crystalline ingots for the photovoltaics (PV) industry is the Czochralski process. It consists of slowly pulling upwards, while simultaneously rotating, an oriented seed out of molten silicon contained in a pure quartz crucible. Float-zone silicon (FZ-Si) is a high-purity and more expensive alternative. In this case, a molten zone is passed along a high-purity polycrystalline rod, heated by a skin current induced by an electromagnetic field, leaving behind a purified monocrystal. There is no direct contact between silicon and the surroundings (except for the inert ambient gas), and the process allows for purification of impurities that segregate in the melt, so contamination is very low. The most widespread method for preparation of mc-Si blocks is the casting process, where the silicon feedstock is melted, cast into molds, and allowed to solidify into blocks. mc-Si is cheaper than mono-Si but has a higher amount of crystal defects and impurities and is also more brittle. However, the improvement in material quality either during the crystallization process or during the cell fabrication has led today to the demonstration of efficiencies as high as 22% on large area [3]. Availability, lower cost, improved wafer quality, and understanding of how to process mc-Si wafers

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into highly efficient solar cells have made them the workhorse of the PV market.

2.1.2 Thin-film silicon In contrast to the fabrication of crystalline silicon wafers, the thin-film silicon technology does not require the costly and delicate crystal traction and sawing. It also avoids kerf loss (about 100 mm per cut wafer), which represents an additional cost factor for c-Si wafers. The technology is based on the deposition of very thin silicon layers on glass or other low-cost substrates at temperatures well below the melting point of c-Si. Although other deposition methods are possible, all types of electronically active silicon films are mostly plasma-deposited from Si- and H-containing precursor gases, typically silane and hydrogen. The most common technique is plasmaenhanced chemical vapor deposition (PECVD), where the substrates are typically mounted on one of the electrodes of a parallel plate structure within a vacuum reactor chamber. Controlled flows of precursor gases are introduced in the chamber, and an RF power is applied between the electrodes. The standard RF frequency is 13.56 MHz, which is the frequency assigned to industrial processes by international authorities. Higher discharge frequencies are also used in the so-called very high frequency PECVD (VHF-PECVD) technique. Compared to crystalline silicon with its well-ordered face-centered cubic lattice, thin-film silicon is typically amorphous (a-Si:H) and contains a substantial amount of hydrogen, incorporated during the deposition. The atomic model structure of a-Si:H is shown in Fig. 2.1A. The silicon atoms largely retain the same basic structure of crystalline silicon, where each atom is linked by covalent bonds to four other silicon atoms arranged as a tetrahedron. However, due to bond angle distortions and dangling bonds, the structure of amorphous silicon presents disorder on medium and long range. As structural uniformity is found only over tiny volumes, by the uncertainty principle the strong spatial localization of charge carriers leaves their momentum largely undetermined. As a result, differently from crystalline silicon (indirect semiconductor), where phonons are required for transitions between valence and conduction band states to satisfy the conservation of momentum, here the requirement on the conservation of momentum is relaxed. The distinction between direct and indirect optical transition disappears and a larger absorption coefficient is found (for the comparison of c-Si and a-Si:H absorption coefficients, see, for example,

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Figure 2.1 (A) Atomic model structure of hydrogenated amorphous silicon. (B) Scheme of the typical microstructure of a mc-Si:H layer with pencil-like conglomerates formed by Si nanocrystals (small white circles) embedded in a-Si:H tissue (gray background).

Ref. [4]). The disorder suppresses the formation of sharp band edges, but a mobility gap between the extended states of the valence- and the conduction-band can still be defined (w1.7 eV). This is accompanied by tails of localized states, which act as shallow trap-states, and by deep defect states close to the middle of the gap, acting as recombination centers. The latter correspond to dangling bonds in the material, whose number needs to be kept as low as possible by passivation with hydrogen. The most intense research about defects in a-Si:H has been focused on light-soaking effects. It was soon discovered that the electronic properties of a-Si:H are degraded by prolonged illumination. This light-induced degradation (LID), known as StaeblereWronski effect, has a self-limiting character, attributed to an increase by a factor of 10 of the dangling bond state density, and is reversible upon annealing [4]. The thin-film silicon family includes also microcrystalline silicon (mc-Si: H), alloys with germanium or carbon, and compounds with oxygen and nitrogen. mc-Si:H consists of small crystallites embedded in an amorphous silicon matrix (Fig. 2.1B). To be precise, it is not one single material, but a class of mixed-phase materials exhibiting a wide range of microstructures depending both on the deposition conditions and on the underlying layer or substrate. In addition, the deposition generally begins with an amorphous phase, and a minimum thickness is needed before a localized phase transformation takes place (nucleation). Once nucleation has occurred, the crystals grow around the seeds. The size of the crystallites is in the 3e10 nm range, and for this reason the material is also known as nanocrystalline silicon (ncSi:H). Like a-Si:H, mc-Si:H can be plasma-deposited from silane and

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hydrogen, but generally with higher power density and very high hydrogen dilution, or in silane-depleted high-pressure regimes. With respect to a-Si: H, mc-Si:H shows no LID or only very moderate, has improved transport properties, and can be more efficiently doped (while doping is self-limited in a-Si:H because the process creates also defects). Its band gap is w1.1 eV as for c-Si, which makes for a perfect combination with w1.7 eV of a-Si:H for integration into full thin-film Si tandem solar cells. More in general, tuning of the band gap is rather easily obtained within the thin-film Si family. Alloying of amorphous Si with Ge allows narrower band gaps down to w1.4e1.5 eV with increasing Ge content (lower values are not feasible because of the poor electronic quality of the material), while band gaps above 1.7 eV are feasible when introducing carbon or oxygen.

2.2 Silicon homojunction solar cells To make a good solar cell, a proper charge extraction scheme is crucial. The photogenerated charge carriers need to be efficiently transported out of the absorber material (the Si wafer in the present case) toward the two external metal terminals while minimizing recombination. Two carrier-selective contacts are needed, which ideally provide for efficient transport of only one type of carrier, while blocking the other type (electrons or holes). There are several ways to implement such electron- and hole-selective membranes on a silicon wafer, giving rise to different types of c-Si solar cells. Here we focus on the homojunction approach (same band gap for the entire structure) going from the classic/industrial to the high-efficiency designs.

2.2.1 Classic design and fabrication process The classic c-Si solar cell (Fig. 2.2A) is based on a moderately p-type doped Si wafer as absorber of light. Doping of the absorber is used to enable efficient transport of majority carriers (holes in this case) to the contacts, but has to be kept low to minimize Auger recombination. The dominant base material at commercial level is p-type silicon with resistivity around 1 Ucm (acceptor density NA w1016 cm3). The primacy of p-type over n-type wafer is essentially due to historical reasons: in the 1950s the application of these “solar batteries” was to power satellites and p-type Si soon became the material of choice, thanks to its better space-radiation hardness. There are also wafer-manufacturing reasons, such as the smaller variation of wafer

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Figure 2.2 (A) Schematic cross-section of the crystalline silicon homojunction solar cell in its classic version and (B) the more advanced design with localized rear contact and dielectric rear surface passivation.

resistivity along the Si ingot plus the relative simplicity and established expertise with phosphorus versus boron diffusion. Within the device, the role of the wafer is to enable efficient absorption of light on a large surface and allow for efficient flow of the photogenerated charge carriers with minimum recombination. The solar cell is then completed by carrier-selective contacts, which are thin highly doped regions at the two surfaces. The electron or hole selectivity is achieved because of the doping-induced shift of the Fermi level toward the conduction or valence band of the absorber, respectively. The large asymmetry between the conductivity for electrons and holes within the highly doped nþ (or pþ) surface regions makes it easy for electrons only (or holes only) to pass through and reach the metal terminals. In the classic design, these selective membranes are obtained by phosphorus and aluminum doping for the front nþ electron collector and the back pþ hole collector, respectively. The solar cell is thus an nþppþ structure, all made of crystalline silicon (homojunction solar cell) with light entering from the nþ side. At the front (nþ region), the donor concentration ND falls steeply from more than 1020 cm3 at the surface to values below NA in a depth of less than 1 mm. At the rear (pþ region), the silicon surface is doped with aluminum within a thickness of w5 mm at a concentration of w3  1018 e 1  1019 cm3. For optical optimization, the wafer surfaces are textured on a length scale of a few micrometers in order to reduce the amount of back reflected light and increase the path length of the photons inside the absorber. Front reflection is further reduced by applying a top antireflection coating based on refractive index matching and quarter-wavelength thickness, with the supplementary role of passivation layer. The front side, exposed to solar light, is

Silicon solar cells: materials, technologies, architectures

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finally completed by a metal grid electrode, while the rear side is fully metalized. The backside metal plays also an optical role, by partially reflecting back into the silicon absorber the long wavelength light not absorbed when first passing through the wafer. The typical fabrication process starts with surface texturing of the wafers by immersion in wet chemical bath, yielding equivalent textures on both the faces. This step also removes the few microns of damaged material resulting from the wire sawing. Anisotropic etching of mono-Si wafers by potassium hydroxide produces random square pyramids, while mc-Si, characterized by grains of many different orientations, requires isotropic etching, usually by acidic solutions. After careful cleaning, the wafers undergo phosphorus diffusion by exposure to phosphorus oxychloride (POCl3) and oxygen in nitrogen carrier within a high-temperature quartz furnace where temperatures of 760e850 C are typically reached. This is the step where the front nþ region (or emitter) is formed together with unwanted equivalently doped regions at the rear and at the edge of the wafer (removed in later steps) due to lack of selectivity of the process. As supplementary beneficial effect, phosphorus diffusion promotes the gettering of metal impurities from the bulk of the wafer, thus making them harmless (this is a key point for mc-Si solar cells). Afterward, the phosphorus silicate glass formed on the wafer during the diffusion step is removed by wet etching and an antireflection coating is applied at the front side (typically hydrogenated amorphous silicon nitride deposited by PECVD). This layer, with refractive index around 2 in the visible range and a thickness around 75 nm for optimal antireflection effect, also plays a passivating role by providing hydrogen termination to the Si dangling bonds at the surface and by field effect due to its high positive fixed charge density. In mc-Si wafers, this layer also allows a beneficial hydrogenation effect on defects within the wafer during the final firing step. The metal contacts are then defined via subsequent printing of different Ag and Al pastes, consisting of metal particles, glass frits as a binder, plus solvents and other organic additives. The front side is printed with H-like patterns consisting of 50- to 100-mm-wide finger gridlines connected by wider busbars. At the rear, pads in the form of a continuous or an interrupted line are firstly printed, followed by full Al coverage. The pads are used to collect the current from the metalized area and to ensure a highly conducting electrode for soldering. Finally, the wafers go through the firing step, which is a sequence of thermal treatments actually forming the front and back contact simultaneously. During this step, the hole-selective contact (or the so-called

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Al-back surface field of the solar cell) is also created because the aluminum forms a eutectic melt with silicon. The unwanted phosphorus-diffused region at the rear side is melted and compensated (overdoped with aluminum). In the cooling phase, silicon recrystallizes with incorporation of aluminum and, at the eutectic temperature, the mixed phase of aluminum and silicon solidifies, thus forming a heavily doped p-type silicon layer. The final step is edge isolation, i.e., the laser ablation of the unwanted phosphorus-diffused layer around the perimeter of the wafer, to prevent parasitic shunting.

2.2.2 High-efficiency designs Several innovative concepts have been proposed to increase the efficiency of c-Si solar cells by addressing the various limitations of the classic technology. The high-efficiency designs have primarily dealt with the high recombination losses at the back side of the classic solar cell, where it is impossible to apply dielectric coatings for passivating the wafer surface, due to the presence of the eutectic mix. The implementation of such insulating layers greatly improves the control of surface recombination. This is the approach of the passivated emitter and rear cell (PERC) family of solar cells [5], where the aluminum-doped pþ region is restricted to a small fraction of the rear surface. Without entering into the specific details, essentially the aluminum paste is screen printed over a rear passivation layer where local openings are previously made so that the pþ regions are formed only within the openings (Fig. 2.2B). Along with the electronic benefits, optical improvements are also achieved, thanks to the enhanced reflectivity at the back side, in the presence of a low refractive index layer between the metal and the silicon wafer, and the suppression of the free-carrier absorption losses with respect to a full-area pþ region. Typically, the rear side of the wafer in these designs is polished in order to fully gain from the integration of the dielectric layer. As for the material choice, silicon oxide formed via high-temperature oxidation is the best option for the passivation, but costwise, deposited films are generally preferred. Aluminum oxide is a good option, with its negative fixed charge that creates a p þ accumulation layer. A few nanometers are deposited, usually capped with thicker silicon nitride for optical optimization. Advances are also achieved with upgraded metallization approaches. The conductivity of the front metal grid can be improved by moving from screen printing to metal plating, allowing for narrower lines and thus reduced

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shading losses. In addition, the firing step is avoided, opening for alternatives for the passivation layers. As for the material choice, copper-based metallization is under investigation to reduce the material costs and circumvent the silver supply issue [6]. Various approaches deal with the front shading losses, like the metal wrap-through (MWT) concept [7]. Looking at the standard solar cell, the large fraction of shading caused by the busbars is evident. MWT solar cells place the busbars on the rear surface. A set of small metalized holes are drilled by laser processing and connect the back busbar through the wafer with the front contact grid. A strategy that pushes this concept even further is the interdigitated back contact (IBC) approach [7]. In this case, both electron- and hole-collecting contacts (the selective membranes and the metal electrodes) are fully placed at the rear side of the solar cell (Fig. 2.3A). This completely eliminates the front metal grid with its shadowing effects and additionally permits a high degree of freedom in the design of the front side to optimize passivation and optics. Such cells are extremely attractive in terms of performance, as they achieve very high short circuit currents, and in terms of aesthetics for their uniform dark look. On the other hand, they pose higher demands on the process technology for the fabrication of the rear surface structure, that is, the separation of p- and n-electrodes. Some of the approaches are the use of diffusion mask layers, ion implantation, and laser doping. To fully exploit the potential of the IBC architecture, n-type wafers are generally preferred (notable exceptions are mentioned in Section 2.5) as they feature higher minority-carrier lifetime in the bulk and superior surface passivation. The textured front side typically presents a silicon oxide passivation layer in combination with silicon nitride as antireflection layer. A

Figure 2.3 (A) Schematic of c-Si solar cells with interdigitated back contact (IBC) design and (B) Si heterojunction solar cell architecture. TCO, transparent conducting oxide

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supplementary lightly diffused layer can improve the passivation by field effect mechanism. At the rear side, boron diffusion is applied to part of the surface to selectively transport holes, while phosphorus is locally diffused in the remaining part to selectively collect electrons. Therefore, the rear side consists of alternating (interdigitated) stripes of p-type and n-type doped regions. For efficient carrier collection, the widths of the stripes are optimized based on the carrier diffusion length and the recombination properties of the phosphorus- and boron-doped regions. The back surface of the wafer is then passivated by a thin insulating layer, and the electrical contact with the doped regions is made through local openings into the passivation layer. Separate metal stripes are aligned over the doped regions, but they are kept slightly narrower to avoid shunting with the neighboring stripes. The metal contacts almost fully cover the rear surface and thus simultaneously act as a back reflector. The IBC solar cell is currently the most complicated and most efficient cSi solar cell in mass production. SunPower has long been in a leading position in the research and development of IBC solar cells. Its top-of-the-line residential solar panels based on this technology now deliver efficiencies up to 22.8% [8]. IBC solar cells are becoming more and more attractive for cell manufacturers, as a promising route for increasing the efficiency for largescale commercial products. Bosch, Samsung, Trina Solar are among the companies involved in the development of this cell type, with the latter recently announcing a record efficiency of 25.04% for large-area cells (243.18 cm2) set in its R&D Labs [9].

2.3 Silicon heterojunction and record solar cells As discussed in the previous section, the Si homojunction solar cell architecture relies on modification of thin regions close to the Si wafer surfaces for selective collection of the photogenerated charge carriers. An alternative archetypal charge extraction scheme consists of the deposition onto the wafer surfaces of doped wider-band-gap thin-films (heterostructure), producing nearly ideal selectivity. As this approach does not make use of thermal diffusion of dopants, there are no major technical limitations to the switch from the more common p-type Si wafer to n-type wafers, which are much easier to passivate and thus hold potential for superior device performance. Besides the use of wider-band-gap materials for the implementation of carrier-selective contacts, there are two key innovations in the Si

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heterojunction (SHJ) solar cell concept. First, it adopts ultrathin hydrogenated amorphous silicon as passivating layer, which allows an atomically sharp interface with the Si wafer when the deposition process is properly tuned. Second, the passivating layer is positioned in between the wafer and the carrier-selective contact. This sequence produces practically recombination-free surfaces, even for the metalized regions, thus explaining the very high open-circuit voltages achieved with this architecture (see the overview at the end of the chapter). From process and cost perspective, major advantages of the SHJ technology are the lower number of steps to fabricate the device, with respect to the high-efficiency silicon homojunction solar cells and the low-temperature (1.7 eV), hydrogenated amorphous silicon oxide is a promising candidate. The development of a material with sufficiently good electronic quality is, however, a challenge since the incorporation of oxygen easily generates defects [19]. A noteworthy approach is the combination of an a-Si:H top cell (EG w 1.75 eV) and a mc-Si:H bottom cell (EG w 1.1 eV) in the socalled micromorph device, where the two materials are an almost perfect match for efficiently sharing the spectral content of the sunlight. The solar cell structure is shown in Fig. 2.5B. With its low band gap, mc-Si:H enlarges the spectral sensitivity of the tandem cell down to the near-infrared region of the AM1.5 solar spectrum, while circumventing the difficulties in achieving high-quality a-SiGe:H with band gap below 1.4 eV. Furthermore, the

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microcrystalline component is more stable than a-Si:H-based alloys. The highest stabilized conversion efficiencies for the thin-film Si technology have been indeed obtained by AIST (Japan) with the combination of amorphous Si and microcrystalline (or nanocrystalline) Si absorbers in tandem cells with two and three components (12.7% and 14.0%, respectively), accompanied by a remarkable 12.3% demonstrated for 1.4 m2 modules by TEL solar Trubbach labs [3].

2.5 Summary and outlook In this chapter, we have briefly reviewed classic and innovative approaches in crystalline and thin-film silicon technologies, of course offering only a partial view onto the wide spectrum of research and development activities due to limited space. Table 2.1 summarizes the performance parameters of the record silicon-based solar cells here discussed, including the achievements based on a selection of the most promising new trends. The thin-film Si technology is a mature and reliable photovoltaic technology with many years of proven outdoor performance. In this approach, the generally expensive semiconductor material is used in minor quantity and can be directly deposited on low-cost large-area substrates. The intrinsically limited efficiency of this technology, however, has put the brakes on market expansion, although with its versatility it remains a valid option for building integration. From the research point of view, many of the achievements are of interest for other fields. One major example is the application of high-quality Si-based thin films to the SHJ solar cells, in the top range of solar cell efficiency. Also in the field of light management, a very advanced level has been reached, providing valuable insight for other technologies. The photovoltaic market is dominated by the crystalline silicon technology with ever-improving efficiency and lowering cost. Several technical developments have been or are being transferred from the research level to the production stage, often readapting costly laboratory solutions while maintaining the key innovations. During 2019, p-type mono-Si PERC is in fact forecast to become the dominant technology deployed by the solar industry for module assembly, with the extraordinary growth in production matched by continuous increase of module power. Thanks to the continuous research, c-Si solar cells are approaching their theoretical Auger-limited efficiency of 29.4% [30], with the present record at 26.7% [3]. Various conceptually different approaches have reached or surpassed the 25% efficiency value, certified even for areas above 100 cm2. As

FF (%)

Efficiency (%)

Single junction

multi-Si p-type (large) Si PERC p-type Si IBC n-type (large) Si TOPCon n-type Si POLO IBC p-type SHJeIBC n-type SHJeIBC n-type (large) a-Si:H (thin film) mc-Si:H (thin film)

Jinko solar UNSW Trina Solar Fraunhofer-ISE ISFH Kaneka Kaneka AIST AIST

245.83 4.00 243.18 4.008 3.9857 79.0 179.74 1.001 1.044

0.6717 0.706 0.7156 0.7241 0.7266 0.738 0.7403 0.896 0.550

40.55 42.7 42.27 42.87 42.62 42.65 42.5 16.36 29.72

80.9 82.8 82.8 83.1 84.3 84.9 84.7 69.8 75.0

22.0 25.0 25.04 25.8 26.1 26.7 26.6 10.2a 11.9a

Multijunction

GaInP/GaAs/Si (4terminal) Perovskite/Si (4terminal) Perovskite/Si (monolithic) a-Si:H/nc-Si:H/nc-Si: H a-Si:H/nc-Si:H (micromorph)

NREL/CSEM/EPFL

1.002

2.52/0.681

13.6/11.0

87.5/78.5

35.9

Imec

0.13b

1.22/0.678

15.4/24.1

73.4/81.2

27.1

Oxford PV

1.030

1.802

19.75

78.7

28.0

AIST

1.045

1.922

9.94

73.4

14.0a

AIST

1.000

1.342

13.45

70.2

12.7a

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IBC, interdigitated back contact; PERC, passivated emitter and rear cell; SHJ, Si heterojunction; TOPCon, tunnel oxide passivated contact; POLO, polysilicon on oxide. a Stabilized by 1000-h exposure to 1-sun light at 50 C. b 0.13 cm2 perovskite cell stacked on top of a 4 cm2 IBC silicon cell in a 4-terminal configuration.

Silicon solar cells: materials, technologies, architectures

Table 2.1 Performance parameters of record archetypal silicon-based solar cells [3,9,10]. Solar cell Institution Area (cm2) Voc (V) Jsc (mA/cm2)

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discussed in this chapter, two elements are required to achieve high efficiency: excellent surface passivation and efficient carrier-selective contacts. Here there is room for novel concepts. Ideally, passivation should also cover the region beneath the metal contacts. Since most passivating materials are insulators, simultaneous suppression of surface recombination and conduction of electrical current is challenging. The use of a semiconductor like intrinsic hydrogenated amorphous silicon as the passivating layer is one excellent option, as in SHJ solar cells. Interesting alternatives under investigation make use of sufficiently thin passivating interlayers to permit quantum mechanical tunneling. The novel tunnel oxide passivated contact (TOPCon) technology developed by Fraunhofer ISE is a significant example that reached a record efficiency of 25.8% with n-type c-Si wafers [3,31]. This approach replaces the traditional screen printed Al-BSF at the rear side of the solar cell with an ultrathin layer of SiO2 (1e2 nm) capped by phosphorus-doped polysilicon layer plus a full silver rear contact. It is remarkable that such high efficiency is reached with a two-side contacted cell. The tunnel oxide concept is under investigation also within the IBC architecture. An efficiency as high as 26.1% has been reported in February 2018 by ISFH employing p-type wafers [3,32]. This cell features a passivating electron-selective nþ-type polysilicon on oxide (POLO) junction at the negative contact of the cell and a passivating hole-selective pþ-type POLO junction at the positive contact, both placed at the rear side. This represents the current world record value for IBC cells on p-type wafer. Alternative dopant-free carrier-selective contacts are also under investigation. Doping presents in fact some drawbacks. In SHJ solar cells, doping of the a-Si:H layers creates defects acting as recombination centers with harmful influence on the passivation. In homojunctions, the heavy doping, required for electrical reasons, enhances Auger recombination. Among the dopant-free concepts, the introduction of transition metal oxide films (in particular molybdenum oxide) as hole-selective contacts in SHJ solar cells, thanks to the high work function, has emerged [33]. The wide band gap of these materials offers the additional advantage of a current density gain. Finally, to really boost the efficiency, huge interest is now placed into the tandem approach, where c-Si can work as a nearly ideal bottom cell absorber in combination with a top cell absorber with band gap around 1.5e1.8 eV, or even more junctions can be considered. In this field, efficiency as high as 35.9% has been demonstrated by the NREL/CSEM/EPFL collaboration for 4-terminal mechanically stacked junctions combining c-Si with IIIeV

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semiconductors with appropriate band gap (GaInP/GaAs/Si) [3,34]. These are interesting results from a performance point of view, while many concerns remain in case of the more elegant 2-terminal monolithic design for the challenges in the heteroepitaxial material growth when considering IIIe V semiconductors. In addition, the costly IIIeV deposition techniques somehow hinder the market penetration of this technology. Meanwhile, organiceinorganic lead halide perovskite solar cells have attracted strong interest due to their potential to increase the efficiency at reasonable cost while being compatible with the c-Si technology. This relatively new entry in the PV sector has quickly gained the scene achieving efficiencies above 20% [3] (see also Chapter 5 in this book). IBC and SHJ architectures are the most promising Si-based bottom counterpart. Experimentally, efficiencies beyond the c-Si single junction record value have been recently demonstrated by Oxford PV and by Imec using 2- and 4-terminal architectures, respectively [3,10,11]. Thanks to the sharp optical absorption edge of perovskites, their tunable band gap in the ideal range for silicon-based tandems, and the low-cost processing, perovskite/silicon tandem solar cells hold great promise for efficiencies above 30% at reasonable production cost soon to come.

References [1] International Technology Roadmap for Photovoltaic (ITRPV), ninth ed., September 2018 available at: http://www.itrpv.net/Reports/Downloads/>. [2] J. Schmidt, R. Peibst, R. Brendel, Surface passivation of crystalline silicon solar cells: present and future, Sol. Energy Mater. Sol. Cell. 187 (2018) 39e54. [3] M.A. Green, E.D. Dunlop, D.H. Levi, J. Hohl-Ebinger, M. Yoshita, A.W.Y. HoBaillie, Solar cell efficiency tables (version 54), Prog. Photovoltaics Res. Appl. 27 (2019) 565e575. [4] A.V. Shah, H. Schade, M. Vanecek, J. Meier, E. Vallat-Sauvain, N. Wyrsch, U. Kroll, C. Droz, J. Bailat, Thin-film silicon solar cell technology, Prog. Photovoltaics Res. Appl. 12 (2004) 113e142. [5] M.A. Green, The passivated emitter and rear cell (PERC): from conception to mass production, Sol. Energy Mater. Sol. Cell. 143 (2015) 190e197. [6] J. Bartsch, A. Mondon, K. Bayer, C. Schetter, M. H€ orteis, S.W. Glunz, Quick determination of copper-metallization long-term impact on silicon solar cells, J. Electrochem. Soc. 157 (10) (2010) H942eH946. [7] E. Van Kerschaver, G. Beaucarne, Back-contact solar cells: a review, Prog. Photovoltaics Res. Appl. 14 (2006) 107e123. [8] https://us.sunpower.com/solar-panels-technology/x-series-solar-panels/(accessed 1.04.2019). [9] https://www.trinasolar.com/us/resources/newsroom/trina-solar-announces-newefficiency-record-2504-large-area-ibc-mono-crystalline (accessed 1.04.2019). [10] M. Jaysankar, B.A.L. Raul, J. Bastos, C. Burgess, C. Weijtens, M. Creatore, T. Aernouts, Y. Kuang, R. Gehlhaar, A. Hadipour, J. Poortmans, Minimizing voltage

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[27] U.W. Paetzold, M. Smeets, M. Meier, K. Bittkau, T. Merdzhanova, V. Smirnov, D. Michaelis, C. Waechter, R. Carius, U. Rau, Disorder improves nanophotonic light trapping in thin-film solar cells, Appl. Phys. Lett. 104 (2014) 131102. [28] K. S€ oderstr€ om, J. Escarré, O. Cubero, F.-J. Haug, S. Perregaux, C. Ballif, UV-nanoimprint lithography technique for the replication of back reflectors for n-i-p thin film silicon solar cells, Prog. Photovoltaics Res. Appl. 19 (2011) 202e210. [29] C. Battaglia, J. Escarré, K. S€ oderstr€ om, M. Charriere, M. Despeisse, F.-J. Haug, C. Ballif, Nanomoulding of transparent zinc oxide electrodes for efficient light trapping in solar cells, Nat. Photon. 5 (2011) 535e538. [30] A. Richter, M. Hermle, S. Glunz, Crystalline silicon solar cells reassessment of the limiting efficiency for crystalline silicon solar cells, IEEE J. Photovolt. 3 (2013) 1184e1191. [31] A. Richter, J. Benick, F. Feldmann, A. Fell, M. Hermle, S.W. Glunz, n-type Si solar cells with passivating electron contact: identifying sources for efficiency limitations by wafer thickness and resistivity variation, Sol. Energy Mater. Sol. Cell. 173 (2017) 96e105. [32] F. Haase, C. Hollemann, S. Sch€afer, A. Merkle, M. Rien€acker, J. Kr€ ugener, R. Brendel, R. Peibst, Laser contact openings for local poly-Si-metal contacts enabling 26.1%-efficient POLO-IBC solar cells, Sol. Energy Mater. Sol. Cell. 186 (2018) 184e193. [33] J. Geissb€ uhler, J. Werner, S. Martin de Nicolas, L. Barraud, A. Hessler-Wyser, M. Despeisse, S. Nicolay, A. Tomasi, B. Niesen, S. De Wolf, C. Ballif, 22.5% efficient silicon heterojunction solar cell with molybdenum oxide hole collector, Appl. Phys. Lett. 107 (2015) 081601. [34] S. Essig, C. Allebé, T. Remo, J.F. Geisz, M.A. Steiner, K. Horowitz, L. Barraud, J.S. Ward, M. Schnabel, A. Descoeudres, D.L. Young, M. Woodhouse, M. Despeisse, C. Ballif, A. Tamboli, Raising the one-sun conversion efficiency of IIIe V/Si solar cells to 32.8% for two junctions and 35.9% for three junctions, Nat. Ener. 2 (2017) 17144.

CHAPTER THREE

Ternary organic solar cells Xiaoling Ma, Fujun Zhang* Key Laboratory of Luminescence and Optical Information, Ministry of Education, Beijing Jiaotong University, Beijing, China *Corresponding author.

3.1 Introduction The continuous increase in fossil fuelebased energy consumption has caused a global environmental problem, and the development of clean and renewable energy resources is an effective approach to solve the energy supply issue in the future [1e3]. Among the numerous renewable energy resources, solar energy is widespread and inexhaustible [4,5]. As a technology of converting sunlight to electricity, organic solar cells (OSCs) have attracted tremendous attention owing to their great potential for large-area production with low-cost materials and simple solution processing method. Currently, most highly efficient OSCs are fabricated with a bulk heterojunction structure in which electron donating and electron accepting materials are blended together within an active layer providing a large interfacial area for exciton dissociation [6,7]. The thickness of active layers is usually very thin (w100 nm) to ensure efficient charge transport and collection. However, the photon harvesting of thin active layers is insufficient due to the intrinsically narrow absorption range of organic materials. Increasing photon harvesting of active layers is the prerequisite issue for obtaining highly efficient OSCs. A detailed investigation on the donor absorption window indicates that the power conversion efficiency (PCE) of OSCs may exhibit approximately 35% improvement if the width of the absorption window can be increased from 200 to 400 nm [8,9]. Compared with the single-junction OSCs, tandem configuration solar cells can not only improve light absorption but also reduce thermalization loss of photonic energy by stacking two or more subcells with complementary absorption range [10e13]. Unfortunately, the tandem OSCs inescapably confront some serious technical challenges due to their complicated multilayers stacking in series; for

Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00003-3

© 2020 Elsevier Ltd. All rights reserved.

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instance, processing the intermediate layers as well as balancing the light absorption among each subcells makes it difficult in practical application. Alternatively, ternary OSCs featuring multiple light harvesting materials in one active layer have emerged as one of the most promising strategies to improve photovoltaic performance since ternary blends possess the advantage of improving photon harvesting ability as in tandem cells, while maintaining the straightforward fabrication that is utilized in single-junction devices [14]. In addition, the incorporation of the third component also can facilitate charge transport, exciton dissociation, and film morphology (molecular crystallinity, crystal orientation, domain size, and purity) [9,15e18]. The key photovoltaic parameters, short circuit current (JSC), open circuit voltage (VOC), and fill factor (FF) can be simultaneously or individually optimized to their maximum values by carefully selecting the third component and adjusting its content. Meanwhile, all the strategies for optimizing the performance of single bulk heterojunction OSCs, such as thermal annealing, solvent vapor annealing, hot solution, mixing solvent, solvent additive, and methanol treatment, can also be effectively applied in ternary OSCs [19e24]. Today, due to rapid development of donor and acceptor materials and device optimization technique, the PCE of ternary OSCs has exceeded 14% [25e27]. In this chapter, the development of ternary OSCs is summarized. In the early studies, fullerene derivatives such as PC61BM ([6,6]-phenyl-C61butyric acid methyl ester) or PC71BM were usually utilized as the dominate electron acceptors in ternary OSCs due to their high electron mobility and isotropy of charge transport. However, fullerene derivatives have several shortcomings, such as weak absorption in the visible region, limited tunability of energy levels, and poor stability [28]. Recently, nonfullerene acceptor materials with the advantages of strong absorptivity in specific region and tunable energy level have attracted much attention of researchers in photovoltaic region. Ternary OSCs can be mainly classified into three categories based on the combination of different acceptor materials: fullerene-based OSCs, fullerene- and nonfullerene-based OSCs, and nonfullerene-based OSCs. A detailed introduction about the fundamental working mechanisms (charge transfer, energy transfer, parallel-linkage, alloy model) of ternary OSCs is presented, followed by corresponding characterization techniques. The study direction of ternary OSCs also involves thick-film device, semitransparent device, and stability for future applications, which is discussed in this chapter. The current challenges and

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further prospects on ternary OSCs are also briefly analyzed in the last section.

3.2 Working mechanism of ternary OSCs It is known that the ternary OSCs can be classified into two categories according to the function of the third component: two donors/one acceptor (D1:D2:A1) and one donor/two acceptors (D1:A1:A2). Nowadays, there are four working mechanisms in both categories of ternary OSCs: charge transfer, energy transfer, parallel-linkage, and alloy model, which is closely related to the location of the third component in ternary active layer. The third component could (1) be fully embedded within one particular host domain; (2) form its own channels; (3) locate at the D/A interface; (4) alloy with the donor or acceptor materials [17], as depicted in Fig. 3.1.

3.2.1 Four working mechanisms In the case of charge transfer mechanism, the incorporation of the third component can provide more additional percolating pathways for efficient exciton dissociation and charge transport. The lowest unoccupied molecular orbital (LUMO) and the highest occupied molecular orbital (HOMO) energy levels of the third component are properly located between the energy levels of the host donor and acceptor to form the cascade energy level alignment. In the D1:D2:A1 ternary system, for example, the third component (D2) acts as a bridge to transfer electrons to acceptor and deliver holes to D1, as shown in Fig. 3.2A. When the third component is located at the

Ternary blend OSCs

Embedded in one phase

Parallel-like-structure

Located at interfaces

Alloy structure

Cathode Electron transport layer Donor + Acceptor + Third component Hole transport layer Anode

Figure 3.1 Structures of ternary blend bulk heterojunction organic solar cells (OSCs) with four possible active layer morphologies. Reproduced with permission L. Lu, M.A. Kelly, W. You, L. Yu, Status and prospects for ternary organic photovoltaics. Nat. Photon. 9 (2015) 491e500. Copyright 2015, Nature.

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Figure 3.2 Schematic of the working mechanisms in ternary organic solar cells based on D1:D2:A1 ternary system. (A) The charge transfer mechanism. (B) The energy transfer mechanism. (C) The parallel-like mechanism. (D) The alloy model. The arrows indicate the possible charge carrier transfer and transport pathway.

interface between donor and acceptor, the holes or electrons generated in the third component can be efficiently collected by the corresponding electrode through the dominating charge carrier transport channels [29,30]. Thus, the observed VOC of ternary OSCs in such systems is basically pinned to the smaller VOC of the original binary OSCs. The energy transfer in ternary blend films should be a competing process compared with the charge transfer among the materials, and only one working mechanism is in a dominant position between energy transfer and charge transfer. In the D1:D2:A1 ternary system (Fig. 3.2B), where the third component only acts as the“energy donor,” all holes are only generated in dominating donor and the third component works just as a sunlight absorber (sensitizer) to extend the absorption window [29,31]. The photoexcited sensitizer can transfer

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energy to host donor or acceptor material via the F€ orster resonance energy transfer (FRET) process [31]. Substantial overlapping of the one material’s emission spectrum and the other material’s absorption spectrum is desirable for an efficient energy transfer. The parallel-linkage mechanism of ternary OSCs is significantly different from the charge transfer and energy transfer mechanisms. In parallel-linkage ternary OSCs, it is believed that the excitons generated in two donors/acceptors can migrate to their respective donor/ acceptor interface and then dissociate into free charge carriers in ternary active layers. In the D1:D2:A1 ternary system, this working mechanism is depicted in Fig. 3.2C. In the parallel-linkage ternary OSCs, the energy transfer and charge transfer are both absent between two donors or two acceptors, equivalent to a parallel connection of two individual binary OSCs. In such ternary systems, the VOC values of ternary OSCs lie between the measured values of the two individual binary OSCs and vary along with the composition of the ternary component, rather than being pinned by the smallest energy level difference between the donors and acceptors. Another working mechanism alloy model with the well-mixed D1/D2 or A1/A2 phases was proposed to explain the tunable VOC in ternary OSCs. It is supposed that the two donors or acceptors electronically couple into a new charge transfer state (CT state), which requires a good miscibility of the two donor or two acceptor materials. The CT state energy in alloy model varies along with the content of ternary blend composition, resulting shifted VOC of ternary OSCs. This working mechanism in a D1:D2:A1 ternary system is depicted in Fig. 3.2D. As discussed above, the four fundamental principles in ternary OSCs (charge transfer, energy transfer, parallel-like, alloy model) are clearly different from each other. The mechanism governs the photovoltaic process in ternary OSCs, and the relevant characterization methods are discussed in the following sections.

3.2.2 Characterization methods In a ternary system, both charge transfer and energy transfer can contribute to the improvement of JSC due to better photon harvesting and more efficient charge generation. Photoluminescence (PL) measurement is a convenient tool to investigate charge transfer or energy transfer between two different materials. Let us refer to the D1:D2:A1 ternary system: if energy transfer exists between two different band gap donors, one would expect increased emission intensity for the relatively low band gap donor with decreased emission intensity for the other donor, considering that the

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quantum yields of the two donors are similar. On the other hand, if charge transfer occurs between two donors, the emission intensity of one donor would be quenched without the increased emission intensity for the other donor. For example, the PL spectra of P3HT:SMPV1 films and solutions were investigated, as shown in Fig. 3.3A and B. It is apparent that the emission intensity of P3HT and SMPV1 are substantially decreased in blend films and solutions along with the incorporation of SMPV1 [32]. This result suggests that charge transfer occur between the two donors. As displayed in

Figure 3.3 Photoluminescence (PL) spectra of P3HT:SMPV1 films (A) and solutions (B) with different SMPV1 content under 490 nm light excitation. (C) PL spectra of P3HT:DIB-SQ films with different DIB-SQ content under 490 nm light excitation. (D) JeV curves of devices with P3HT, SMPV1, or P3HT:SMPV1 (1:1) as active layers. (E) JeV curves of devices with P3HT, DIB-SQ, or P3HT:DIB-SQ as active layer. (A) and (B) Reproduced with permission of American Chemical Society from Q. An, F. Zhang, L. Li, J. Wang, Q. Sun, J. Zhang, W. Tang; Z. Deng, Simultaneous improvement in short circuit current, open circuit voltage, and fill factor of polymer solar cells through ternary strategy. ACS Appl. Mater. Interfaces 7 (2015) 3691e3698. (C) Reproduced with permission of American Chemical Society from Q. An, F. Zhang, L. Li, J. Wang, J. Zhang, L. Zhou, W. Tang, Improved efficiency of bulk heterojunction polymer solar cells by doping low-bandgap small molecules. ACS Appl. Mater. Interfaces 6 (2014) 6537e6544. (D) Reproduced with permission of American Chemical Society from Q. An, F. Zhang, L. Li, J. Wang, Q. Sun, J. Zhang, W. Tang; Z. Deng, Simultaneous improvement in short circuit current, open circuit voltage, and fill factor of polymer solar cells through ternary strategy. ACS Appl. Mater. Interfaces 7 (2015) 3691e3698. (E) Reproduced with permission of American Chemical Society from Q. An, F. Zhang, L. Li, J. Wang, J. Zhang, L. Zhou, W. Tang, Improved efficiency of bulk heterojunction polymer solar cells by doping low-bandgap small molecules. ACS Appl. Mater. Interfaces 6 (2014) 6537e6544.

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Fig. 3.3C, the emission intensity of P3HT is gradually reduced while DIBSQ exhibits continuously improved emission intensity along with the increase of DIB-SQ content in P3HT:DIB-SQ blend films under 490 nm light excitation, suggesting efficient energy transfer from P3HT to SQ molecules [33]. It should be noticed that there are some possible errors with the PL measurement of film, induced by the sample thickness, measure angle, or molecular alignment, and so on. Zhang et al. first proposed a simple method, named as pure donor/acceptor device, to verify the charge transfer or energy transfer between two donors or acceptors [32,33]. These devices should be fabricated with only donors or acceptors as active layers. The JSC of P3HT: SMPV1-based devices is higher than that of P3HT- or SMPV1-based devices (Fig. 3.3D), which should be ascribed to efficient charge transfer between two donors. The JSC of P3HT:DIB-SQ-based devices is in between those of P3HT-based and DIB-SQ-based devices (Fig. 3.3E), suggesting that charge transfer between P3HT and DIB-SQ should be negligible. Yu et al. also fabricated a solar cell using only two donors to prove the charge transfer at the donors interfaces, exhibiting a relatively low PCE of 0.051% with the corresponding low JSC ¼ 0.192 mA cm2, VOC ¼ 0.65 V, and FF ¼ 40.8% [18]. Combining PL measurement and pure donor/acceptor device is an efficient and exact strategy to distinguish charge transfer and energy transfer between two different materials, which has been demonstrated in many systems by different groups [34,35]. F€ orster theory predicts a decrease in the excited-state lifetime of the FRET-donor with increasing FRET-acceptor concentrations, as FRET introduces an additional nonradiative decay path for the donor into the system. Taylor and coworkers studied the energy transfer between P3HT and SQ by femtosecond fluorescence upconversion technique [36]. The lifetime of neat P3HT film is 223 ps, which is decreased to 52.4 ps by incorporating 1% SQ and further to 9.9 ps with 5% SQ, suggesting the existence of energy transfer from P3HT to SQ. In PTB7-Th:BTR:PC71BM ternary system, the time-resolved transient photoluminescence spectra of the neat PTB7-Th, BTR, and their blend films were investigated by monitoring 760 nm light emission corresponding to the PTB7-Th emission peak [37]. The emission lifetime is increased from 0.08 to 0.18 ns along with the increased content of BTR, which further demonstrates the occurrence of FRET from BTR to PTB7-Th. The efficiency of FRET is given by: sDA s¼1  (3.1) sD

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where sD and sDA are the fluorescence lifetime of the FRET-donor in the absence and presence of the FRET-acceptor, respectively [38]. In P3HT: SQ:PC71BM system, the energy transfer efficiency is 77% or 96% for the blend films with 1% or 5% SQ [36]. The similar study of energy transfer efficiency was also carried out by Zhang’s group in SMPV1:DIB-SQ: PC71BM system. The energy transfer efficiency was about 56% or 67% for blend films SMPV1:SQ with weight ratio of 9:1 or 1:1, respectively [39]. Time-resolved transient absorption (TA) spectrum is another method to monitor the photophysical processes. Taylor and coworkers further investigated the photophysical processes of P3HT and SQ by TA spectra, as shown in Fig. 3.4A. TA spectra usually compose two main features: ground state bleach (GSB) with negative signature and excited state absorption (ESA) with positive signature. Neat P3HT film and blend P3HT:SQ films were excited at 500 nm to selectively excite P3HT. At early time, the

Figure 3.4 (A) Transient absorption (TA) spectra of neat P3HT film and P3HT:SQ blend films. The pumping wavelength was 500 nm with a fluence of 8 mJ cm2. The picosecond TA spectra of (B) neat PBDB-T, (C) neat PTB7-Th, and (D) blend PBDB-T:PTB7Th (80:20, wt/wt) films. (E) Transient absorption dynamics of neat PBDB-T, PTB7-Th, and their blend films by probing 585 and 720 nm, respectively. (A) Reproduced with permission of Springer Nature from J.S. Huang, T. Goh, X. Li, M.Y. Sfeir, E.A. Bielinski, S. Tomasulo, M.L. Lee, N. Hazari, A.D. Taylor, Polymer bulk heterojunction solar cells €rster resonance energy transfer. Nat. Photon. 7 (2013) 479e485. (E) Reproemploying fo duced with permission of Wiley-VCH from X. Ma, Y. Mi, F. Zhang, Q. An, M. Zhang, Z. Hu, X. Liu, J. Zhang, W. Tang, Efficient ternary polymer solar cells with two well-compatible donors and one ultranarrow bandgap nonfullerene acceptor. Adv. Energy Mater. 8 (2018) 1702854.

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negative signal contains three peaks at 520, 560, and 610 nm which well accord with the P3HT absorption with 0e2, 0e1, and 0e0 vibrational transitions, respectively. The positive signal at 650 nm represents the photoinduced ESA of P3HT. The TA signal is changed by introducing 1% or 5% SQ. The negative signal represents the GSB of SQ, suggesting that SQ is excited. The excitation of SQ is caused by the excited P3HT rather than the pumping wavelength, due to the SQ negligible absorption at 500 nm. Meanwhile, signals of the photoinduced absorption and stimulated emission in P3HT both substantially decay to zero. This suggests that the excitation energy is transfer from P3HT to SQ. The similar phenomenon can also be observed from PBDB-T:PTB7-Th:IEICO-4F-based OSCs [40], as shown in Fig. 3.4BeE. The TA spectra of neat PBDB-T and PTB7-Th films exhibit featured GSB and ESA signal. For the PBDB-T:PTB7-Th blend films, the GSB and ESA features are the overlapped features of two donors. The ESA peak of PTB7-Th in blend films shows an enhanced intensity and slow decay z1 ps compared to that in the neat PTB7-Th film. The energy transfer process from PBDB-T to PTB7-Th should occur due to the decreased PBDB-T signature and the simultaneously enhanced PTB7-Th signature. The lifetime of GSB signal by probing 585 nm is decreased from 0.49 to 0.47 ps by incorporating 20 wt% PTB7-Th. The lifetime of GSB signal by probing 720 nm is 4.5 ps for the neat PBT7-Th and 6.1 ps for PBDB-T:PTB7-Th blend films, respectively. The energy transfer from PBDB-T to PTB7-Th can be further confirmed from the enhanced lifetime of PTB7-Th GSB signal (720 nm) in the blend films. The efficient energy transfer can enhance the exciton utilization and contribute to PCE improvement. The concept of parallel-linkage was first proposed by You and coworkers, in which ternary blends can be regarded as two independent binary subcells [41]. Two groups of materials, TAZ:DTBT:PCBM (0.5:0.5:1 weight ratio) and DTffBT:DTPyT:PCBM (0.5:0.5:1 weight ratio), were employed for constructing parallel-linkage ternary OSCs. To keep the same quantity of individual components in corresponding ternary and binary OSCs, the thickness of binary and ternary active layers are fixed at w50 and w100 nm, respectively. The w50-nm-thick binary OSCs are defined as “subcell” for better understanding. The absorption spectra of parallel-linkage ternary OSCs are essentially linear combinations of the spectra of their two“subcells,” as shown in Fig. 3.5A and B. The EQE spectra of binary OSCs and ternary OSCs were also measured and are shown in Fig. 3.5C and D. The EQE spectra of parallel-linkage ternary

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(A) 0.6

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Figure 3.5 Absorption spectra of the parallel-linkage ternary organic solar cells (OSCs) and their “subcells” based on (A) TAZ/DTBT and (B) DTffBT/DTPyT. EQE spectra of the parallel-linkage ternary OSCs and their “subcells” based on (C) TAZ/DTBT and (D) DTffBT/DTPyT. JeV curves of the parallel-linkage ternary OSCs and their “subcells” based on (E) TAZ/DTBT and (F) DTffBT/DTPyT. Reproduced with permission of American Chemical Society from L. Yang, H. Zhou, S.C. Price, Y. Wei, Parallel-like bulk heterojunction polymer solar cells. J. Am. Chem. Soc. 134 (2012) 5432e5435.

OSCs are approximately the sum of those of the individual “subcells” in the short wavelength range, where both donor polymers contribute to the formation of electronehole pairs. This suggests that most of the free charge

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carriers generated in each “subcell” of the PBHJ device are successfully collected by the respective electrodes. In the long-wavelength range, the EQE spectra of the parallel-linkage ternary OSCs are higher than that of the smaller band gap “subcell.” This is probably due to that the large band gap polymer with high mobility can serve as an additional charge transport channel in the parallel-linkage ternary OSCs, which facilitate the charge transport. In general, the JSC of the parallel-linkage ternary OSCs are significantly increased and almost identical to the sum of those for the two single “subcells”; the VOC values of the parallel-linkage ternary OSCs are between those of individual “subcells,” as shown in Fig. 3.5E and F. This should be attributed to that holes generated from individual donor polymers mainly travel through their corresponding polymer-connected channel to the anode, which is similar to the parallel connection of two single-junction binary OSCs. The parallel-linkage mechanism in ternary OSCs can also be verified from the PL spectra of the corresponding neat films or blend films. In PTB7:PBDT-TS1:PC71BM ternary system [42], both PL emission peaks of PTB7 and PBDT-TS1 are at about 780 nm due to their similar band gap of 1.8 eV. The PL emission intensity of PBDT-TS1 films is stronger than that of PTB7 films under the same excitation conditions. The emission intensity of PTB7:PBDT-TS1 blend films is gradually enhanced in the range from 750 to 850 nm along with the increased content of PBDT-TS1; and the emission peak of the blend films is located at between that of neat PTB7 and PBDT-TS1 film. This indicates neither charge transfer nor energy transfer exist between PTB7 and PBDT-TS molecules. In addition to D1:D2:A1 ternary system, the parallel-linkage mechanism has also been demonstrated between two acceptors. Sun et al. reported a ternary system based on PDBT-T1:PCBM:ITIC-Th as active layer [35]. The parallellinkage can be characterized by grazing incidence X-ray diffraction (GIXD) and resonant soft X-ray scattering (RSoXS) technology. When the ITIC-Th content increased to more than 30 wt%, the diffraction peak of PDBT-T1, PCBM, and ITIC-Th domain can be observed, suggesting that the used three materials can form their individual phase. For RSoXS measurements, 0% ITIC-Th blend shows a shoulder in the scattering at w0.014Å1, corresponding to a length scale of 45 nm for the phase separated domain. When the ITIC-Th content increased up to 50 wt%, the RSoXS profile shows two shoulders in the scattering at 0.013 Å-1 (48 nm) and 0.003 Å-1 (210 nm), indicating the formation of multilength scale morphology. The two length scales should be derived from the ITIC-Th

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and PCBM-rich domains. As a result, the PDBT-T1:PCBM:ITIC-Thbased OSCs work in parallel-linkage structure model. The alloy model (with the well-mixed D1/D2 or A1/A2 phases) was proposed by Thompson and Street et al. [43] to explain the tunable VOC of ternary OSCs. In this mechanism, the two donors or two acceptors should be well mixed to form the same frontier orbital energy levels, which are determined by blend composition. The VOC of OSCs is determined by the energy of interface band gap (EGI) minus the quasi-Fermi energies, where the EGI is the energy difference value between the HOMO energy level of donor and the LUMO energy level of acceptor. The reason of continuous change in VOC is the corresponding change of HOMO or LUMO energy levels. The photocurrent spectral response (PSR) was carried out by Thompson et al. to explore the electronic states in ternary blends and confirm the continuous change in the energy of the HOMO or LUMO of the complementary two-component materials. The PSR of the mixed acceptor system P3HT:PCBM:ICBA was measured and are shown in Fig. 3.6A. It has been reported that the energy below the optical band gap of each material is derived from the direct excitation at interface from the HOMO of donor to the LUMO of acceptor (Fig. 3.6A inset). Previous studies have shown that the actual interface band gap corresponds to a (B) 1.5

100 Fraction ICBA

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Figure 3.6 (A) Photocurrent spectral response (PSR) data for the P3HT:PCBM:ICBA ternary system plotted as a function of ICBA fraction in PCBM:ICBA pair. The inset represents the EGI or charge transfer transition. (B) Plot of the estimated interface band gap energy defined by PSR at photocurrent values of 0.1 and 0.01 compared to the values of VOC for P3HT:PCBM:ICBA ternary system. Reproduced with permission of American Chemical Society from R.A. Street, D. Davies, P.P. Khlyabich, B. Burkhart, B.C. Thompson, Origin of the tunable open-circuit voltage in ternary blend bulk heterojunction organic solar cells. J. Am. Chem. Soc. 135 (2013) 986e989.

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photocurrent of w104 of the peak value, where the shape of the PSR spectrum changes from an exponential at lower energy to a broader band at higher energy absorption of the heterojunction interface [44]. The interface absorption measured at photocurrent values of 0.1 and 0.01 (dashed lines in Fig. 3.6A) and the VOC of ternary OSCs were summarized in Fig. 3.6B. It can be observed that VOC is w0.55 V smaller than the interface band gap, which is consistent with the expected quasi-Fermi energies. The results indicate that the change in the EGI was accurately measured by VOC and that EGI changes continuously with the composition of two-component acceptors. The composition dependence of the band gap for an alloy model is often described by the extension of Vegard’s law [45,46]. EG1 ¼ ð1  XÞEG1 þ XEG1  bxð1  xÞ

(3.2)

in which b is known as the bowing factor and EG1 and EG2 are the band gaps of the used two materials. This model fits the data well, which can be observed in Fig. 3.6B. Similar measurements are also carried out in other ternary system based on P3HT75-co-EHT25/P3HTT-DPP-10%/PCBM as active layers. In this system, the VOC also can be continuously changed along with the different proportions of two donors. Either two donors or two acceptors can form an electronic alloy with changed HOMO and LUMO energy levels based on the average composition of two donors or acceptors, resulting in variation of VOC with composition in these two components. The alloy model can also be verified by varied HOMO and LUMO energy level of donors or acceptors, which can be calculated by cyclic voltammetry (CV) curves. In PBDB-T:ITCPTC:IDT6CN-M-based system [47], the LUMO energy levels of the acceptor blend films are gradually decreased along with the increase of ITCPTC content, indicating that the alloy state of ITCPTC and IDT6CN-M is formed in the acceptor blend films, as shown in Fig. 3.7A and B. In another ternary system based on PSTZ:IDIC:ITIC as active layer [48], the ITIC:IDIC blends exhibit the varied frontier orbital energies dependence on the composition of these two components, which indicate the forming of an electronic alloy between ITIC and IDIC, as shown in Fig. 3.7C. Recently, another strategy was proposed to distinguish if there exists an alloy model between two acceptors. The linearly varied VOC can be observed in ternary system when the two acceptors have similar chemical structure. A quasilinear correlation of VOC could be observed in the ternary system where the chemical nature of the two acceptors is drastically different. If the alloy states can be formed between two donors or

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(B) Ferrocene

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Figure 3.7 (A) Cyclic voltammetry plots of ITCPTC:IDT6CN-M blend films with different ITCPTC content. (B) The HOMO and LUMO energy levels of the acceptor blend films against the ITCPTC content in the acceptors. (C) Electronic energy levels of the ITIC: IDIC blend films with different ITIC contents. (D) Dependence of the VOC of the organic solar cells on the MeIC1 content, energy levels of the used materials, and schematic dynamic processes in the ternary active layers. (E) Pictorial representations of the frontier molecular orbitals of INPIC-4F and MeIC1 from the DFT calculations. Diagram of the frontier orbital energies of INPIC-4F and MeIC1. LUMO, lowest unoccupied molecular orbital; HOMO, highest occupied molecular orbital. (A) and (B) Reproduced with permission of Royal Society of Chemistry from M. Zhang, W. Gao, F. Zhang, Y. Mi, W. Wang, Q. An, J. Wang, X. Ma, J. Miao, Z. Hu, Efficient ternary non-fullerene polymer solar cells with PCE of 11.92% and ff of 76.5%. Energy Environ. Sci. 11 (2018) 841e849. (C) Reproduced with permission of Elsevier from W. Su, Q. Fan, X. Guo, X. Meng, Z. Bi, W. Ma, M. Zhang, Y. Li, Two compatible nonfullerene acceptors with similar structures as alloy for efficient ternary polymer solar cells. Nano Energy 38 (2017) 510e517. (E) Reproduced with permission of Royal Society of Chemistry from X.L. Ma, W. Gao, J.S. Yu, Q.S. An, M. Zhang, Z.H. Hu, J.X. Wang, W.H. Tang, C.L. Yang, F.J. Zhang, Ternary nonfullerene polymer solar cells with efficiency > 13.7% by integrating the advantages of the materials and two binary cells. Energy Environ. Sci. 11 (2018) 2134e2141.

two acceptors, the mixed new quasi frontier orbitals density (Ne) of alloy states strongly depends on the content of the third component. The VOCs of ternary OSCs can be calculated by the equation: VOCTernary ¼

f 1  N e1  V OCBinary1 þ f 2  N e2  V OCBinary2 f 1  N e1 þ f 2  N e2

(3.3)

in which V OCBinary1 , V OCBinary2 , and VOCTernary are the VOCs of corresponding binary OSCs and ternary OSCs; Ne1 and Ne2 are the Ne of acceptor 1 (donor

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1) and acceptor 2 (donor 2); f1 and f2 are the weight ratios of acceptor 1 (donor 1) and acceptor 2 (donor 2) in acceptors (donors). Using DFT quantum calculations, Ne can be obtained by Ne ¼ nl, where n is the molecular number of unit mass, and l is the number of quasi-degenerate LUMOs ( 11% are summarized in Table 3.1. It can be observed that most of the highly efficient ternary OSCs are based on nonfullerene acceptor or the combination of fullerene and nonfullerene acceptors. The combination of ternary strategy with-high performance material will continually show great potential in future development.

3.4 The potential research directions of ternary OSCs In addition to PCE improvement, there are also some potential research directions of ternary OSCs for future application, such as thickfilm, semitransparent device, and stability, which will be discussed in this section.

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PCE [%]

References

PCE10:IDTBR p-DTS(FBTTh2)2:PC71BM PSTZ:ITIC J52:IT-M PffBT4T-2OD:PC71BM PTB7-Th:meta-TrBRCN PTB7-Th:PC71BM PBDB-T:ITIC FTAZ:IDTC J71:ITIC J71:IT-M PBDB-T:IEICO-4F PBDB-T:IDT6CN-M PTB7-Th:COi8DFIC PTB7-Th:BDTThIT-4F PTB7-Th:PC71BM PBDB-T:IT-M PBDTTT-EFT:IEICO-4F PBDB-T:IT-M PBDB-T:IT-M PBDB-T:SFBRCN PTB7-Th:F8IC PBT1-C:MeIC PBDB-T:NNBDT PTB7-Th:ITIC-2F PTB7-Th:COi8DFIC PBT1-C:ITIC-2Cl PBDB-T:IT-M PTB7-Th:3TT-FIC BTR:NITI PBDB-T:INPIC-4F PM6:ITCPTC PBDB-T-2F:IT-4Cl PTB7-Th:COi8DFIC

11.0 11.0 11.1 11.1 11.17 11.4 11.4 11.41 11.6 11.6 11.6 11.62 11.92 11.94 12.03 12.10 12.10 12.15 12.16 12.20 12.27 12.3 12.55 12.8 12.9 13.08 13.4 13.52 13.54 13.63 13.73 14.13 14.18 14.62

[103] [104] [48] [105] [106] [107] [108] [109] [110] [111] [85] [40] [47] [112] [113] [114] [109] [115] [116] [117] [118] [119] [88] [120] [121] [122] [123] [124] [125] [82] [51] [89] [25] [27]

d 8.77 8.06 9.4 10.57 10.15 9.21 10.03 10.4 10.7 10.68 10.25 10.51 10.72 9.61 10.10 11.10 10.63 10.89 10.80 9.39 10.9 11.47 11.7 12.1 11.47 11.1 11.71 12.21 9.03 12.55 13.04 13.45 10.48

IDFBR ZnP IDIC IEICO PCDTBT8 ITIC-Th BTR N2200 ITIC-Th-O MeIC2 ITIC PTB7-Th ITCPTC IEICO-4F IEICO-4F DR3TSBDT N2200 PCDTBT ITCN Bis-PC70BM PTB7-Th PC71BM MeIC2 FDNCTF IOTIC-2F BDTThIT-4F ICBA P1 PC71BM PC71BM MeIC1 MeIC IT-2Cl PC71BM

PCE, power conversion efficiency; OSCs, organic solar cells.

3.4.1 Thick active layerebased ternary OSCs It is known that ternary OSCs have attracted much attention due to that ternary organic blends can broaden the absorption range of OSCs without the use of complicated tandem cell structures. Despite their broadened

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absorption range, the light harvesting capability of ternary OSCs is still limited because ternary OSCs mostly use thin active layers of about 100 nm in thickness, which is not sufficient to absorb all photons in their spectral range and may also cause problems for future roll-to-roll mass production that requires thick active layers. Huang et al. reported a very efficient ternary OSC by incorporating a nematic liquid crystalline small molecule (BTR) into a PTB7-Th:PC71BM host system [108]. Compared with binary systems, a higher PCE of 10.37% is achieved for ternary OSCs with 25 wt% BTR content at the 250-nm-thick active layer. From GIWAXS characterization of blend films (Fig. 3.15A), the PTB7-Th (010) peak relating to pep stacking in the out-of-plane direction shifts to a higher q from 1.53 Å1 to 1.67 Å1 along with BTR content gradually increased to the optimal ratio (25%), and then shifts back with the further incorporation of BTR. This suggests that the blend film with 25% BTR has the most compact pep stacking with d-spacing of around 3.77Å. The more ordered molecular arrangement should facilitate hole transport, which can be confirmed from measured hole mobility according to the space charge limited current method. The ternary OSCs have a shorter charge extraction time than that of the binary devices with the same thick active layers (Fig. 3.15B), which is consistent with the change of charge mobility. Recently, Huang et al. reported an all-polymer OSC with PBTA-Si:PTzBI-Si:N2200 as ternary active layer [126]. The optimized all-polymer OSCs attain a PCE of 9.17% with an active layer thickness of 350 nm and maintain a PCE over 8% for thicknesses over 400 nm. Fig. 3.15C illustrates the dark JeV curves for device A (PBTA-Si:N2200based OSCs), device B (PTzBI-Si:N2200-based OSCs), and device C (PBTA-Si:PTzBI-Si:N2200-based OSCs); the thicknesses of active layer for those device is appropriately 300 nm. By fitting the dark JeV curves with a one diode replacement circuit [127], the diode ideality factor (n) was obtained, i.e., 1.86, 1.74, and 1.76 for devices A, B, and C, respectively. A lower ideality factor indicates weaker recombination induced by defect states. Although device C has an ideality factor similar to that of device B, the series resistance (Rs) of device C (1.46 U cm2) is lower than that of device B (Rs ¼ 1.74 U cm2) based on the JeV characteristics in the dark. Using the equivalent circuit model [128], the FF-containing series and parallel resistance losses were calculated, and the values of 69.37%, 73.85%, and 75.47% are obtained for devices A, B, and C, respectively. This result implies that device C, i.e., the optimized ternary OSC has a better potential to be fabricated and used in thick-film devices. The ternary strategy should be

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Figure 3.15 (A) In-plane and out-of-plane GIWAXS profiles for PTB7-Th:BTR:PC71BM blend films with various contents of BTR and 100 nm thickness. (B) Transient photocurrent of organic solar cells with PTB7-Th:BTR:PC71BM ratios of 1:0:1.1 and 0.75:0.25:1.1 with thin and thick active layers. (C) Dark JeV curves and the corresponding ideality factors for devices AeC with thicknesses of z300 nm (fitted by a one diode replacement circuit). (B) Reproduced with permission of American Chemical Society from G. Zhang, K. Zhang, Q. Yin, X. F. Jiang, Z. Wang, J. Xin, W. Ma, H. Yan, F. Huang, Y. Cao, Highperformance ternary organic solar cell enabled by a thick active layer containing a liquid crystalline small molecule donor. J. Am. Chem. Soc. 139 (2017) 2387e2395. (C) Reproduced with permission of Wiley-VCH from B. Fan, P. Zhu, J. Xin, N. Li, Y. Lei, W. Zhong, Z. Li, W. Ma, F. Huang, Y. Cao, High-performance thick-film all-polymer solar cells created via ternary blending of a novel wide-bandgap electron-donating copolymer. Adv. Energy Mater. 8 (2018) 1703085.

an efficient way to achieve high-performance thick-film devices that can meet the needs of future roll-to-roll production.

3.4.2 Semitransparent ternary OSCs Semitransparent OSCs have recently attracted extensive attention due to their potential applications in building integrated photovoltaics, wearable electronic power, and photovoltaic vehicles [129e131]. Ternary strategy can also be applied in semitransparent device to improve PCE and maintain

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average visible transmittance (AVT) of semitransparent OSCs. Sun et al. [132] fabricated a ternary semitransparent OSC based on PTB7-Th: PBT1-S:PC71BM as ternary active layer. The PCE of semitransparent OSCs is improved from 8.0% to 9.2%, which should be ascribed to enhanced light absorption and optimized film morphology. More importantly, the addition of a small amount of PBT1-S to the binary blend did not significantly change its AVT and neutral color perception, as shown in Fig. 3.16A and B. In addition to fullerene-based semitransparent system, ternary strategy was first applied in nonfullerene system by Zhang et al. [113]. The development of ultranarrow band gap nonfullerene materials provides more opportunity for the fabrication of semitransparent device. A semitransparent ternary OSC was fabricated with PTB7-Th:BDTThIT4F as host system and an ultranarrow band gap material IEICO-4F as the third component. The PCE of optimized semitransparent ternary OSCs is improved to 9.40% with AVT simultaneously improved to 24.6%, which should be among the highest values for semitransparent OSCs. The PCE improvement should be mainly attributed to enhanced photon harvesting in long wavelength region, which can be confirmed from the EQE spectra shown in Fig. 3.16C. The transmittance of Au/Ag electrode is relatively high in the spectral range from 300 to 600 nm, which is beneficial to prepare semitransparent OSCs. The AVT of semitransparent OSCs is increased from 24.2% to 24.6%, which can be calculated from the transmission spectra of corresponding OSCs shown in Fig. 3.16D. The ultranarrow band gap material IEICO-4F can also be used as third component in other ternary system with one narrow band gap material PTB7-Th as donor and two ultranarrow band gap materials (COi8DFIC and IEICO-4F) as acceptors [112]. The PCE of semitransparent OSCs is increased to 8.23% with simultaneously improved AVT of 20.78% by incorporating 15 wt% IEICO-4F in acceptors. Unlike the commonly reported role of the third component in improving the performance of ternary OSCs, IEICO-4F can induce a variation in the COi8DFIC molecular arrangement [122], resulting in extended photon harvesting in the NIR region, which can be confirmed from the EQE spectra of OSCs (Fig. 3.16E). To qualitatively show the transparency of semitransparent OSCs for human eyes, the photograph of a peony flower partly covered with semitransparent OSCs is exhibited in Fig. 3.16F. It is apparent that the peony flower covered with semitransparent OSCs can be clearly observed, suggesting the potential of semitransparent OSCs in the building-integrated photovoltaic field. In Zhang’s work [112,113], the PCE and AVT of semitransparent OSCs can be simultaneously improved

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Figure 3.16 (A) Optical transmittance of PTB7-Th:PC71BM, PBT1-S:PC71BM-, PTB7-Th: PBT1-S:PC71BM-based semitransparent organic solar cells (OSCs) covered by different Ag thickness (14 mm  15 nm). (B) The representation of color coordinates of PTB7Th:PC71BM, PBT1-S:PC71BM-, PTB7-Th:PBT1-S:PC71BM-based semitransparent OSCs under AM 1.5G illumination on the CIE chromaticity diagram. (C) EQE spectra of blended PTB7-Th:BDTThIT-4F:IEICO-4F films with different IEICO-4F contents. (D) Transmission spectra of corresponding films and semitransparent OSCs. (E) EQE spectra of PTB7Th:COi8DFIC:IEICO-4F-based semitransparent OSCs with different IEICO-4F contents. (F) Flower partly covered by semitransparent OSCs with (1) PTB7-Th:COi8DFIC; (2) the optimized ternary system; (3) PTB7-Th:IEICO-4F as the active layers. (B) Reproduced with permission of Wiley-VCH from Y. Xie, L. Huo, B. Fan, H. Fu, Y. Cai, L. Zhang, Z. Li, Y. Wang, W. Ma, Y. Chen, Y. Sun, High-performance semitransparent ternary organic solar cells. Adv. Funct. Mater. 28 (2018) 1800627. (D) Reproduced with permission of Elsevier from Z. Hu, J. Wang, Z. Wang, W. Gao, Q. An, M. Zhang, X. Ma, J. Wang, J. Miao, C. Yang, F. Zhang, Semitransparent ternary nonfullerene polymer solar cells exhibiting 9.40% efficiency and 24.6% average visible transmittance. Nano Energy 55 (2019) 424e432. (F) Reproduced with permission of Royal Society of Chemistry from X. Ma, Z. Xiao, Q. An, M. Zhang, Z. Hu, J. Wang, L. Ding, F. Zhang, Simultaneously improved efficiency and average visible transmittance of semitransparent polymer solar cells with two ultra-narrow bandgap nonfullerene acceptors. J. Mater. Chem. 6 (2018) 21485e21492.

by employing ternary strategy. Ternary strategy may provide more opportunity for fabricating efficient semitransparent OSCs.

3.4.3 Stability of ternary OSCs Many results confirm that appropriate incorporation of the third component is beneficial to improve device stability. Baran et al. fabricated a ternary OSC

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based on P3HT:IDTBR:IDFBR as active layer, and the PCE is improved from 6.3% to 7.7% with 30 wt% IDFBR in acceptors [103]. Combined with another low-band gap donor polymer PTB7-Th, the two mixed acceptors also lead to the ternary OSCs a higher PCE of 11.0%. The ambient stability of ternary and binary OSCs was tested. After 1200 h in air and under dark conditions (Fig. 3.17A), the ternary P3HT:IDTBR:IDFBR (1:0.7:0.3) device retains 80% of its PCE (6.1%), whereas P3HT:IDTBR performance retains 70% (4.3%). In addition, the P3HT:IDTBR, P3HT:IDFBR, and P3HT:IDTBR:IDFBR devices were exposed at operating conditions (unencapsulated, in air, AM1.5 radiation to illumination 100 mW cm2) for an initial 90 h test (Fig. 3.17B). The P3HT:IDTBR:IDFBR device exhibits the best air photostability, retaining 85% of its initial performance after 90 h. These results suggest that the addition of IDFBR to P3HT:IDTBR blend not only improves photovoltaic performance but also has a synergistic benefit on both storage lifetime and photostability. The stability of ternary OSCs was also investigated by Zhang et al. In nonfullerene-based ternary system PBDB-T:ITIC:N2200, [109] the optimized ternary OSCs exhibit excellent stability with about 80% initial PCE after 1000 h storage; however, only 74% or 42% initial PCEs can be retained for the corresponding binary OSCs after 1000 h storage (Fig. 3.17C). It means that the ternary active layers with 10 wt% N2200 exhibit a more stable morphology compared with other blend films, which can be confirmed from the GIXD (Fig. 3.17D) and TEM (Fig. 3.17E) experimental results. In fullerene-based ternary system PTB7-Th:BTR:PC71BM, the stability was investigated in inverted and conventional OSCs [37]. The decay trends of other photovoltaic parameters dependent on the storage time were fitted according to the power function (Y ¼ a(1 þ T)b), in which Y represents the key photovoltaic parameter, T represents the storage time, the coefficient a is the initial value of the corresponding photovoltaic parameter, and the absolute value of index (jbj) reflects the degradation rate. The smaller jbj values of the ternary OSCs indicate that the stability of the OSCs can be improved by employing the ternary strategy. The preserved percent of PCE of the inverted OSCs is 77.0% and 81.8% for the PTB7-Th:PC71BM-based OSCs and the optimized ternary OSCs, respectively. The preserved percent of PCE of the conventional OSCs is 65.5% and 70.2% for the corresponding cells. Obviously, the stability of the OSCs can be enhanced by employing BTR as a nucleating agent to form a more stable morphology.

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Figure 3.17 (A) Shelf storage lifetime (dark, in air) comparison of P3HT:IDTBR:IDFBR device efficiencies with other polymer:fullerene systems. Devices were exposed to ambient conditions over 1200 h duration or until high devices no longer showed any diode behavior. (B) Photostability of P3HT:IDTBR:IDFBR device and polymer:fullerene solar cells (in air, unencapsulated, under AM1.5 illumination at 1 sun) for 90 h. (C) Stability of PBDB-T:ITIC-, PBDB-T:ITIC:N2200-, PBDB-T:N2200-based OSCs without encapsulation under AM 1.5G illumination at intensity of 100 mW cm2. (D) Out-of-plane (solid line) and in-plane (dotted line) GIXD profiles of blend films with different N2200 content. (E) TEM images of blend films with different N2200 content. (B) Reproduced with permission of Springer Nature from D. Baran, R.S. Ashraf, D.A. Hanifi, M. Abdelsamie, N. €hr, S. Holliday, A. Wadsworth, S. Lockett, M. Neophytou, Reducing the Gasparini, J. A. Ro efficiency-stability-cost gap of organic photovoltaics with highly efficient and stable small molecule acceptor ternary solar cells. Nat. Mater. 16 (2017) 363e369. (E) Reproduced with permission of Elsevier from Q. An, F. Zhang, W. Gao, Q. Sun, M. Zhang, C. Yang, J. Zhang, High-efficiency and air stable fullerene-free ternary organic solar cells. Nano Energy 45 (2017) 177e183.

3.5 Challenges and outlooks As discussed in previous section, ternary strategy should be a simple, effective, and promising approach to realize performance improvement of OSCs. It should be noticed that the emerging of high-performance nonfullerene acceptors provides flexible options to rationally design and construct high-performance ternary OSCs. Results reported so far reveal that the

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fullerene:nonfullerene or nonfullerene materialsebased ternary OSCs should be the most promising ternary system, both of which have achieved a PCE more than 14%. In addition to increased photon harvesting, an appropriate incorporation of the third component can offer multiple benefits and synergistic effects, including optimizing film morphology, molecular arrangement, and enhancing charge carrier mobility as well as promoting exciton dissociation at the interface, thus resulting in performance improvement of OSCs. There are, however, still some challenges and limitations for the development of ternary OSCs: (1) compared with binary OSCs, the incorporation of the third component brings more complexities and variables into the binary system, and a deep understanding of the working mechanism in ternary OSCs is required for the further performance improvement. Advanced detection techniques such as TA spectroscopy should be employed to unravel the underling photophysical mechanism in ternary systems. (2) Deeper understanding of the VOC behavior for ternary systems is necessary. It was found that the VOC of ternary OSCs is sensitive to the film morphology. In alloy model, the VOC in ternary OSCs varies as a function of composition, while in some not particularly compatible systems, the VOC of ternary OSCs is mostly determined by the smallest energy level difference between donor and acceptor components. In some rare cases, the VOC in ternary OSCs is even higher than that of two binary OSCs. It is necessary to further investigate the origin of VOC in the ternary system and establish the relationship between morphology and device physics. (3) It is difficult to unify the guidelines of selecting the third component that matches well with the host materials. In addition to material properties (absorption spectrum, energy levels alignment, crystallinity, and surface energy) of the third component, the photovoltaic parameters of two binary OSCs should also be taken into consideration to construct effective ternary OSCs. In terms of those factors, different system may match different selection criteria, which makes the selection criteria ill-defined. (4) For donor materials, previous investigations mainly focused on developing low band gap polymer due to the poor light absorbing ability of fullerene acceptors. Nowadays, most of highperformance nonfullerene acceptors show strong absorption in long wavelength region (700e950 nm). Therefore, high-performance medium band gap or wide band gap donor materials should be further developed to match nonfullerene acceptors’ energy level and complement their absorption band. For future applications, more investigation on thick active layerebased ternary OSCs, the mechanisms of improved stability in ternary OSCs, and

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ternary semitransparent OSCs should be carried out. The PCE of 15% is commonly regarded as the target for enabling commercial viability. It is believed that the ternary OSCs based on fullerene:nonfullerene or nonfullerene materials as acceptors would be the most promising ternary system. The combination of high-performance materials and ternary strategy provide a bright prospect to push the performance of OSCs to the stage of commercial applications.

Acknowledgments This work was supported by National Natural Science Foundation of China (61675017, 61564003, 61377029).

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[85] Z.H. Hu, F.J. Zhang, Q.S. An, M. Zhang, X.L. Ma, J.X. Wang, J. Zhang, J. Wang, Ternary nonfullerene polymer solar cells with a power conversion efficiency of 11.6% by inheriting the advantages of binary cells, ACS Energy Lett. 3 (2018) 555e561. [86] M. Zhang, W. Gao, F.J. Zhang, Y. Mi, W.B. Wang, Q.S. An, J. Wang, X.L. Ma, J.L. Miao, Z.H. Hu, X.F. Liu, J. Zhang, C.L. Yang, Efficient ternary non-fullerene polymer solar cells with pce of 11.92% and ff of 76.5%þ, Energy Environ. Sci. 11 (2018) 841e849. [87] J.X. Wang, W. Gao, Q.S. An, M. Zhang, X.L. Ma, Z.H. Hu, J. Zhang, C.L. Yang, F.J. Zhang, Ternary non-fullerene polymer solar cells with an efficiency of 11.6% by simultaneously optimizing photon harvesting and phase separation, J. Mater. Chem. 6 (2018) 11751e11758. [88] Q. An, J. Zhang, W. Gao, F. Qi, M. Zhang, X. Ma, C. Yang, L. Huo, F. Zhang, Efficient ternary organic solar cells with two Compatible non-fullerene materials as one alloyed acceptor, Small 14 (2018) 1802983. [89] T. Liu, Z. Luo, Q. Fan, G. Zhang, L. Zhang, W. Gao, X. Guo, W. Ma, M. Zhang, C. Yang, Y. Li, H. Yan, Use of two structurally similar small molecular acceptors enabling ternary organic solar cells with high efficiencies and fill factors, Energy Environ. Sci. 11 (2018) 3275e3282. [90] Z. Zheng, Q. Hu, S. Zhang, D. Zhang, J. Wang, S. Xie, R. Wang, Y. Qin, W. Li, L. Hong, A highly efficient non-fullerene organic solar cell with a fill factor over 0.80 enabled by a fine-tuned hole-transporting layer, Adv. Mater. (2018) 1801801. [91] S. Zhang, Y. Qin, J. Zhu, J. Hou, Over 14% efficiency in polymer solar cells enabled by a chlorinated polymer donor, Adv. Mater. 30 (2018) 1800868. [92] J. Sun, X. Ma, Z. Zhang, J. Yu, J. Zhou, X. Yin, L. Yang, R. Geng, R. Zhu, F. Zhang, Dithieno[3,2-b:2’,3’-d]pyrrol fused nonfullerene acceptors enabling over 13% efficiency for organic solar cells, Adv. Mater. 30 (2018) 1707150. [93] P.Q. Bi, F. Zheng, X.Y. Yang, M.S. Niu, L. Feng, W. Qin, X.T. Hao, Dual forster resonance energy transfer effects in non-fullerene ternary organic solar cells with the third component embedded in the donor and acceptor, J. Mater. Chem. 5 (2017) 12120e12130. [94] X. Xu, Z. Bi, W. Ma, Z. Wang, W.C.H. Choy, W. Wu, G. Zhang, Y. Li, Q. Peng, Highly efficient ternary-blend polymer solar cells enabled by a nonfullerene acceptor and two polymer donors with a broad composition tolerance, Adv. Mater. 29 (2017) 1704271. [95] L. Gao, Z.G. Zhang, L. Xue, J. Min, J. Zhang, Z. Wei, Y. Li, All-polymer solar cells based on absorption-complementary polymer donor and acceptor with high power conversion efficiency of 8.27%, Adv. Mater. 28 (2016) 1884e1890. [96] Y. Guo, Y. Li, O. Awartani, H. Han, J. Zhao, H. Ade, Y. He, D. Zhao, Improved performance of all-polymer solar cells enabled by naphthodiperylenetetraimide-based polymer acceptor, Adv. Mater. 29 (2017) 1700309. [97] T. Kim, J.H. Kim, T.E. Kang, C. Lee, H. Kang, M. Shin, C. Wang, B. Ma, U. Jeong, T.S. Kim, Flexible, highly efficient all-polymer solar cells, Nat. Commun. 6 (2015) 8547. [98] H. Benten, T. Nishida, D. Mori, H.J. Xu, H. Ohkita, S. Ito, High-performance ternary blend all-polymer solar cells with complementary absorption bands from visible to near-infrared wavelengths, Energy Environ. Sci. 9 (2016) 135e140. [99] W.Y. Su, Q.P. Fan, X. Guo, B. Guo, W.B. Li, Y.D. Zhang, M.J. Zhang, Y.F. Li, Efficient ternary blend all-polymer solar cells with a polythiophene derivative as a hole-cascade material, J. Mater. Chem. 4 (2016) 14752e14760. [100] S.C. Price, A.C. Stuart, L.Q. Yang, H.X. Zhou, W. You, Fluorine substituted conjugated polymer of medium band gap yields 7% efficiency in polymer-fullerene solar cells, J. Am. Chem. Soc. 133 (2011) 4625e4631.

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[101] K. Li, Z.J. Li, K. Feng, X.P. Xu, L.Y. Wang, Q. Peng, Development of large bandgap conjugated copolymers for efficient regular single and tandem organic solar cells, J. Am. Chem. Soc. 135 (2013) 13549e13557. [102] Z.Y. Li, L. Ying, R.H. Xie, P. Zhu, N. Li, W.K. Zhong, F. Huang, Y. Cao, Designing ternary blend all-polymer solar cells with an efficiency of over 10% and a fill factor of 78%, Nano Energy 51 (2018) 434e441. [103] D. Baran, R.S. Ashraf, D.A. Hanifi, M. Abdelsamie, N. Gasparini, J.A. R€ ohr, S. Holliday, A. Wadsworth, S. Lockett, M. Neophytou, Reducing the efficiencystability-cost gap of organic photovoltaics with highly efficient and stable small molecule acceptor ternary solar cells, Nat. Mater. 16 (2017) 363e369. [104] N. Li, G. Ke, Y. Jiang, Q. Rong, X. Hu, Y. Dong, L. Feng, X. Peng, T.P. Russell, G. Zhou, Small-molecule solar cells with simultaneously enhanced shortcircuit current and fill factor to achieve 11% efficiency, Adv. Mater. 29 (2017) 1700616. [105] R. Yu, S. Zhang, H. Yao, B. Guo, S. Li, H. Zhang, M. Zhang, J. Hou, Two wellmiscible acceptors work as one for efficient fullerene-free organic solar cells, Adv. Mater. 29 (2017) 1700437. [106] W. Li, J. Cai, F. Cai, Y. Yan, H. Yi, R.S. Gurney, D. Liu, A. Iraqi, T. Wang, Achieving over 11% power conversion efficiency in pffbt4t-2od-based ternary polymer solar cells with enhanced open-circuit-voltage and suppressed charge recombination, Nano Energy 44 (2018) 155e163. [107] W. Wu, G. Zhang, X. Xu, S. Wang, Y. Li, Q. Peng, Wide bandgap molecular acceptors with a truxene core for efficient nonfullerene polymer solar cells: linkage position on molecular configuration and photovoltaic properties, Adv. Funct. Mater. 28 (2018) 1707493. [108] G. Zhang, K. Zhang, Q. Yin, X.F. Jiang, Z. Wang, J. Xin, W. Ma, H. Yan, F. Huang, Y. Cao, High-performance ternary organic solar cell enabled by a thick active layer containing a liquid crystalline small molecule donor, J. Am. Chem. Soc. 139 (2017) 2387e2395. [109] Q. An, F. Zhang, W. Gao, Q. Sun, M. Zhang, C. Yang, J. Zhang, High-efficiency and air stable fullerene-free ternary organic solar cells, Nano Energy 45 (2017) 177e183. [110] P. Cheng, J. Wang, Q. Zhang, W. Huang, J. Zhu, R. Wang, S.Y. Chang, P. Sun, L. Meng, H. Zhao, Unique energy alignments of a ternary material system toward high-performance organic photovoltaics, Adv. Mater. 19 (2018) 1801501. [111] J. Wang, W. Gao, Q. An, M. Zhang, X. Ma, Z. Hu, J. Zhang, C. Yang, F. Zhang, Ternary non-fullerene polymer solar cells with efficiency of 11.6% by simultaneously optimizing photon harvesting and phase separation, J. Mater. Chem. 6 (2018) 11751e11758. [112] X. Ma, Z. Xiao, Q. An, M. Zhang, Z. Hu, J. Wang, L. Ding, F. Zhang, Simultaneously improved efficiency and average visible transmittance of semitransparent polymer solar cells with two ultra-narrow bandgap nonfullerene acceptors, J. Mater. Chem. 6 (2018) 21485e21492. [113] Z. Hu, J. Wang, Z. Wang, W. Gao, Q. An, M. Zhang, X. Ma, J. Wang, J. Miao, C. Yang, F. Zhang, Semitransparent ternary nonfullerene polymer solar cells exhibiting 9.40% efficiency and 24.6% average visible transmittance, Nano Energy 55 (2019) 424e432. [114] T. Kumari, S.M. Lee, S.H. Kang, S. Chen, C. Yang, Ternary solar cells with a mixed face-on and edge-on orientation enable an unprecedented efficiency of 12.1%, Energy Environ. Sci. 10 (2017) 258e265. [115] M.J. Xiao, K. Zhang, S. Dong, Q.W. Yin, Z.X. Liu, L.Q. Liu, F. Huang, Y. Cao, High-performance ternary nonfullerene polymer solar cells with both improved

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[131] Y. Cui, C. Yang, H. Yao, J. Zhu, Y. Wang, G. Jia, F. Gao, J. Hou, Efficient semitransparent organic solar cells with tunable color enabled by an ultralow-bandgap nonfullerene acceptor, Adv. Mater. 29 (2017) 1703080. [132] Y. Xie, L. Huo, B. Fan, H. Fu, Y. Cai, L. Zhang, Z. Li, Y. Wang, W. Ma, Y. Chen, Y. Sun, High-performance semitransparent ternary organic solar cells, Adv. Funct. Mater. 28 (2018) 1800627.

CHAPTER FOUR

Dye-sensitized solar cells: from synthetic dyes to natural pigments Antonino Bartolotta, Giuseppe Calogero CNR-IPCF, Istituto per i Processi Chimico-Fisici, Messina, Italy

4.1 Introduction The energy conversion from one form to another has been one of the foremost challenges to promote the worldwide economic, technological, and political progress. New advanced highly efficient technologies to obtain electrical power have become a primary need, due also to the drastic growing worldwide population and the impressive territorial urbanization. Nowadays the primary energy sources are derived from the carbon-based fossil fuels which will fall short over the long term. In addition, their continuous combustion produces huge emissions of greenhouse gas (GHG), leading to serious climate changes and harmful side effects to human health. To address this problem the renewable energies, especially solar energy, have stimulated the scientific and industrial interest. Solar energy is an inexhaustible available source which can be directly converted into electrical power without any environment affect [1,2]. The sun is our primary carbonneutral energy source which strikes the earth with w4.3∙1020 J/hdan energy amount quite equivalent to that consumed on the planet in a year (4.1∙1020 J) [3]. Approaches to capture and exploit energy from the sun could provide benefits well beyond any better expectation. Overcoming the huge gap between our present use of photovoltaic (PV) energy and its enormous potential is the challenge in the coming years. The PV effect takes place on the circumstance that photons (hn) falling on a semiconductor can produce electronehole (e-/hþ) pairs and that at a junction between two different materials, this effect gives rise to an electrical potential difference DV (volt ¼ joules/coulomb) across the interface. In this

Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00004-5

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framework, both inorganic and electrochemical PV solar cells (PVCs) operate upon the establishment of an electric potential difference between two electronically dissimilar materials (n- and p-type) or between an n- or ptype semiconductor and a redox electrolyte. Harvesting energy from the sun can be accomplished via inorganic (e.g., mono- or poly-crystalline silicon, CdTe, CdS, CIGS [4]) and organic PV technologies [5] (e.g., blends of fullerene derivatives (phenyl-C61-butyric acid methyl-ester (PCBM)) as well as, via conjugated polymers (poly-3-hexil-thiophene (P3HT)) [6] as typical) or via photoelectrochemical devices (e.g., dye sensitized solar cells, DSSCs [7,8]) based on hybrid structures of inorganic semiconductors and molecular assemblies, the aim being always to obtain the maximum conversion power at the minimum cost. Obviously, different PV materials will have different energy band gaps (Eg) and thus different light absorption ability. As a general rule, photons with energy equal to the material Eg will be absorbed generating free e- while, photons with lower energy than the material Eg will pass through this latter. On the other hand, photons with higher energy than the Eg will release their excess energy in the form of heat (see Fig. 4.1A). So far, the market PVCs has been dominated by doped forms of single- or poly-crystalline silicon [9] (see Fig. 4.2A), and energy conversion efficiencies (h) exceeding 25% have been achieved in crystalline-silicon heterojunction solar cell [10]. With the prospect to optimize materials usage and manufacturing, maintaining high efficiencies and costing saving, a second generation of single- and multijunction thin films technology based on amorphous silicon [11], cadmium-telluride (CdTe) [12], or cadmium-telluride/ cadmium-sulfide (CdTe/CdS) [13], copper indium gallium di-selenide (CIGS) [14e16]) is becoming a continuously expansive area. However, despite significant developments over the past decades, the energy cost associated to extraction/recycling/purification of limited availability elements (i.e., gallium, germanium, indium, tellurium) might cause some raw material constraints, which in turn will affect the cost production. For instance, purification to metallurgic-grade Si has high energy demand and creates significant amounts of hazardous by-products [17]. On the contrary, its less expensive amorphous counterpart (a-Si) suffers the StaeblereWronski degradation [18], capable to collapse the device’s h up to 30%, already in the first months of use [19]. Also, potential toxicity and related environmental risks associates to specific elements (e.g., Cd is a highly toxic heavy metal material [20]) pose further criticality to large-scale sustainability of these technologies.

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Figure 4.1 (A) Schematic description of energy level diagram of semiconductor band gap; (B) typical structure of a solar cell based on a “pen junction” between two inorganic semiconductors characterized by free negative electrons (n-type) and free positive holes (p-type), respectively. Such kind of assembly allows the photogenerated electrons and holes to be separated and transferred to external circuit for electrical energy production. (C) Typical scheme of full organic solar cell.

At the beginning of 20th century, ecological evaluations relative to the toxic or GHG emission including global warming have become the driving force in promoting alternative energy source. Fig. 4.2B and C shows the amount of GHGs emitted by the diverse electricity sources and the increase of renewable energy sources demand in the world energy consumption, respectively. Although the PV technology cannot be envisaged as the major contribution to the primary energy need without ultralow-cost options for massive distribution and advances in the material and physic science; however, DSSCs have been the subject of a considerable research interests (see Fig. 4.2D) due to their ideal compromise between efficiency and cost performances [21]. Despite less attractive for large-scale implementations, where highly efficient but more expensive single-crystal Si and IIIeV multijunction PV devices are more practicable, DSSCs apply a different mechanism offering several advantages compared to the more traditional PVCs. In fact, while in conventional PV systems the semiconductor assumes both the tasks of light absorption and charge carrier transport, in DSSCs the two functions are separated. Light is absorbed by a sensitizer, which is anchored to the surface of a wide band

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Figure 4.2 (A) The market share of the most common inorganic photovoltaic (PV) materials. (B) Comparison of CO2 emission of electricity generation sources. Greenhouse gas emission by solar PV is a small fraction than that emitted by lignite. (C) Plot of the world energy demand for the main fossil fuels (coal, oil, and natural gas) compared to the shift in the request for renewable energy sources (hydrothermal, wind, biomass, and solar) until 2030). (D) Number of publications per year from literature search using the keywords “dye sensitized solar cells.” (C) Adapted by M. G. O’Regan, A low-cost, highefficiency solar cell based on dye-sensitized colloidal TiO2 films, Nature 353 (1991) 737e740; H.J. Snaith, Estimating the maximum attainable efficiency in dye sensitized solar cells. Adv. Funct. Mater. 20 (2010) 13e19. Data source: ISI Web of Knowledge.

semiconductor. Charge separation takes place at the interface via photoinduced e- injection from the sensitizer into the semiconductor’s conduction band (CB). In addition, while in conventional PVCs decreasing the incoming light intensity (diffusive light) decreases the charge carrier mobility mc (cm2/ V$s), thus declines the cell performance, DSSCs show stable performances at no-standard temperature conditions, irradiation and solar incident angle. Moreover, DSSCs can be manufactured using roll-to-roll processing [22,23] without vacuum or high temperatures. All these favorable features, together with easy integration into building architecture and the possibility to realize DSSCs with on demand semitransparency, multicolor options, etc. [24e26], make them the closely and appealing choice to mimic the natural photosynthesis in the solar-to-electricity conversion.

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Following the pioneering Gr€atzel’s cell in early 1990 [27], and the theoretically predicted conversion efficiency h, of w20% [28], several industrial players have launched campaigns to facilitate DSSCs’ commercial availability and, to enhance their efficiency, stability, and integration into modules and panels. In 2009, G24 Power Ltd. (G24i) [29] first commercialized DSSC products for indoor and portable electricity applications. 3GSolar Photovoltaics Ltd. [30], manufactured effective DSSC modules able to work even under occluded and low-angle light and, developed miniature DSSCs (from 2 cm2 to >200 cm2) for wireless electronics with device’s lifetimes more than 10 years and reliable power output under indoor and/or low-light conditions. Improvements in different areas, including electrodes [31e33], electrolytes [34], photosensitizers [35] as well as the overall device planning, in order to develop efficient DSSCs, have been reached. In 2014, Mathew et al. [36] achieved an h of 13% in a DSSC based on a molecularly engineered porphyrin dye (SM315) and a Co3þ/2þ redox mediator. Reliable values of VOC (0.91 V), JSC (18.1 mA cm2), and FF (0.78) were also reported. The incorporation of new materials such as graphene (Gr) and related materials with their exceptional electrical, optical, and mechanical properties, has been extensively exploited as part of each DSSC components: transparent electrodes to replace FTO at the photoanode [37], light absorbers [38], additives to improve charge transport through both the TiO2 and the electrolyte [39,40], materials for Pt replacing at the CE [41,42]. In 2015, Kakiage et al. [43] reached h exceeding 14% in DSSCs containing Au/graphene nanoplatelets (GrNPs) as a counter electrode, Co3þ/2þ as redox couple, and an LEG4þADEKA-1 as a sensitizer. In 2017, Sudhakar et al. [44] reported an interesting study where, an optimized Gr amount dispersed in o-dichlorobenzene and uniformly incorporated in a suitable polymer was spin coated over the cell photoanode as a sort of barrier to limit the back electrons transfer processes. The fabricated DSSC showed an overall h of 10.4%, which was w28% higher than that of reference cell employing conventional TiO2-based photoanode.

4.2 Dye-sensitized solar cells: structure and operating principles DSSCs provide a technically and economically credible alternative concept to current-day pen junction PV devices, with breakthroughs in the field of physic, chemistry, materials science, and nanotechnology,

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offering new opportunities in indoor and portable energy applications [7,45,46]. DSSCs (also called Gr€atzel’s cell) depart completely from the classical solid-state junction device, by replacing the contacting phase to the semiconductor by a liquid, gel, or solid electrolyte, thus forming a photoelectrochemical device offering attractive features of cost content and environmental advantages. At the heart of the system is the photoanode consisting of a monolayer of sensitizer (dye) adsorbed onto a semiconductor mesoporous oxide layer (w10 mm thick) of nanometer-sized particles (usually TiO2-anatase, although ZnO, SnO2, Nb2O5, etc. have also been exploited) permeated by an electrolyte solution containing a redox couple (commonly I-3/I) acting as donor electron to oxidized dye molecules. Under light irradiation, the dye promotes an e- from its highest occupied molecular orbital (HOMO) to the lowest unoccupied molecular orbital (LUMO) whose energy level is higher than the semiconductor’s CB edge into the CB of the oxide. The excited electrons are injected into the semiconductor’s CB from where they diffuse to be collected to a transparent conductive oxide (fluorine-doped tin oxide (FTO), or indium tin oxide (ITO)) and then, passing through the external circuit, reenter the cell at the counter electrode (CE, typically a Pt-sputtered/TCO glass) to reduce the “hole transport media” (e.g., I-3 / I). Finally, the reducing agent (e.g., I ions) diffusing from the CE into the semiconductor film regenerates the oxidized dye (ms time range) to its original state and completing the oxidation/reduction cycle. In such a pathway the cell generates electrical power from light without suffering any chemical transformation. The sequential events in a DSSC, when I/I-3 is used with TiO2, can be summarized as follows [45]:

- Dye þ hy / Dye*LUMO - Dye*LUMO þ TiO2 / Dyeþ þ eCB-(TiO2) - 2Dyeþ þ 3I /2 Dye þ I3 - 2ecatalyst þ I3 / 3I - Dye* / Dye - eCB-(TiO2) þ DyeþHOMO / Dye þ TiO2 - 2 eCB- (TiO2) þ I3 / 3I þ TiO2

e photoexcitation from HOMO to LUMO levels of the dye e injection in the TiO2 CB Dye regeneration Electrolyte regeneration (reduction) Decay of the Dye excited state e migration from TiO2 CB to dye HOMO level (back reaction) e migration from TiO2 CB to electrolyte (dark current)

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The last three reactions representing the undesirable loss pathways. Optimization of DSSC requires that both electron injection and dye regeneration pathways by electron donor mediator, must be faster than the decay of Dye* and the back reaction of the injected electron with the oxidized dye, respectively. The net electrons flow into the external circuit will result by the kinetic competition between the charge carriers transport and the recombination processes between injected electrons and redox mediator and/or oxidized sensitizer. In this respect, remarkable characteristic of DSSCs, responsible also of their relatively high conversion efficiency (h), is the low recombination kinetics between photoinjected electrons and I-3 ions [7,47] as well as the negligible recombination (ms to ms timescale) with oxidized dye (at high electron density or light intensity [7,48]) compared to the electron injection dynamics (fs to ps time scale). Other fundamental requirement for both efficient charge injection into the semiconductor’s CB and oxidized dye regeneration by the redox couple is that the dye LUMO and HOMO [49] energy levels, the former, sufficiently high in energy for an efficient charge injection and, the second sufficiently low for efficient oxidized dye regeneration. In dark conditions, the semiconductor’s quasi-Fermi level (EF) equals the redox potential Eredox of the redox mediator and no net current flows. Under illumination, the semiconductor’s EF shifts up as the electron concentration increases and a driving force for the electrons to perform electrical work is obtained. The maximum cell voltage output VOC (usually z 08 O 0.9 V using I/I-3 couple) will be determined by the difference between the quasi-Fermi level EF of the electron in the semiconductor’s CB (0.5 to - 0.4 eV vs. NHE for TiO2) and the redox potential of the electrolyte mediator (- 0.4 eV vs. NHE using I/I-3 couple). In this regard, the shift of the redox mediator electrochemical potential toward more positive values and/or the shift of the dye ground state potential toward less positive values would allow to improve the cell VOC. Fig. 4.3 shows a schematic cross section of a conventional DSSC, along with a rough depiction of the individual component energy levels and characteristic times for the different charge transfer steps.

4.2.1 Analytic model and photovoltaic performance 4.2.1.1 Diode equivalent circuit model The modeling optimization of any PV device passes necessarily by the understanding of the physical mechanism underlying its working behavior. According to this thought, mathematical models [50,51], and experimental

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Figure 4.3 (A) Simplified design of a conventional dye-sensitized solar cell. (B) Approximate energy diagram and dynamics of the processes involved in a dye-sensitized nanocrystalline TiO2 solar cell. The favorable pathways (1: electron injection, 2: charge collection, 3: dye regeneration) determining the best cell performance (current and/or voltage), as well as those which thwart the successful operation (4: luminescence or nonradiative decay, 5e6: recombination) are shown with red and blue arrows, respectively. (C) Dynamics of redox processes involved in the light-to-electric power conversion. NHE refers to the standard hydrogen electrode reference. The two common strategies to improve the cell h: (i) shift of the dye’s ground state potential to a less positive value (maintaining the excited-state potential constant) to allow a higher light collection and greater JSC and/or (ii) shift of the redox shuttle’s electrochemical potential to a more positive value, to assist a VOC enhancement, are also indicated.

works [52,53], to cognize the better physical configuration and the electric features of the individual elements, have been reported. Among these models, often differing in the number of parameters intricate in the cell voltage and current calculations, the classical one diode equivalent circuit model (1DEC) has been recognized as a powerful tool to illustrate the interior physical and chemistry properties of the device. In fact, although characterized by a rather high level of abstraction, 1DEC allows to evaluate the DSSC electrical mechanisms (i.e., voltage and current) and to design the whole system with a minimal combination of electrical components (e.g., resistors, diodes, and current sources), and avoid the use of a large number of physiochemical parameters. To make clearer to the reader, we will briefly introduce the main features of this simple model capable to simulate the cell’s network and links its PV characteristic with the internal properties of the individual components. Fig. 4.4A shows the 1DEC diagram for a conventional DSSC where, a series resistance (RS, U∙cm2) is in series with a diode (representing the electrical properties of

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Figure 4.4 (A) Classical one-diode-based equivalent electrical circuit for dye-sensitized solar cells. (B) Influence of Rsh and Rs values on the IeV curve slop. (C) Currentevoltage curve of a photovoltaic cell in the dark (red line) and under illumination (blue line). Under dark condition there is almost no current flowing, until a threshold value at which the contacts start to inject heavily at forward bias for voltages larger than the open circuit voltage. Differently, under illumination the maxima generated photocurrent flows at short-circuit current condition ① and, at ② the photocurrent is balanced to zero (flat band condition). Between ① and ② the cell generates power with a maximum value at the point ③, where the product between current and voltage is the largest. The ratio between the generated power output and the incident solar power determines the overall cell efficiency.

the pen junction) and with a shunt resistance (RSh, U∙cm2), while a constant current (Iph) is in parallel with the junction. The relation between external cell current and voltage circuit can be written as: 
qðV I,Rs Þ V  I$RS I ¼ Iph  I0 e m,kB T  1 þ (4.1) RSh where V is the applied voltage, while Rs takes into account the losses due to macroscopic resistance contributions (i.e., bulk transport, interface transfer, and transport through the contacts), and Rsh comprises the back-electron transfer across the semiconductor/dye/electrolyte junction. Also, the

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diode and Iph represent the electron transfer at the semiconductor/ dye/electrolyte interface and the photogenerated current (equal to the sum of the current through the diode and the shunt resistance), respectively, and I0 the dark current (diode saturation current). Finally, m is the diode’s ideal factor, able of giving information about the dominant recombination mechanism, q and kBT being the elementary charge and the thermal energy, respectively. For highly efficient DSSCs, m values ranging from 2.2 up to 2.7 have been reported [54,55]. The extension of 1DEC model via the addition of more diodes and resistors permits to describe more complicated solar cell systems (e.g., DSSC modules) as well as to include the contribution of supplementary loss mechanisms (recombination) [56e58]. Giannuzzi et al. [59] reported a modified model, which including a second diode to capture the behavior of the reverse bias on a DSSC module (made by 4 W connected cells) perfectly parallels the current/voltage experimental data. The resistive effects dissipating power across parasitic internal resistances reduce the cell efficiency with a heavy impact on the fill factor (FF), the parameter that provides practical information about the cell maximum power output (Pmax). In order to achieve high FF, very large RSh (theoretically /N), to prevent leakage currents and very low RS (theoretically / 0), to get a sharp rise in the forward current, are preferred. In a rough manner, it is possible to derive RS and RSh from the slope of currentevoltage curve at V ¼ VOC and at I ¼ ISC, respectively. The characteristic influence of both resistive contributions is shown in Fig. 4.4B. Also, a typical IeV curve, measured by varying from zero (short circuit) to infinite (open circuit) an external load, under standard 1 full sun condition (AM 1.5-type simulated sunlight irradiation [60]), is presented in Fig. 4.4C. 4.2.1.2 Evaluation of DSSC performance The overall energy light/electrical power conversion efficiency of DSSC can be expressed as [45]: h ¼

JSC $VOC $FF Pmax ¼ Pinc Pinc

(4.2)

where JSC (mA∙cm2) is the short circuit current density, i.e., the photogenerated current per unit active area of the photoanode when the voltage across the cells is zero (short circuit condition); VOC (V) is the open circuit voltage, correlated, as already mentioned, to the difference between the

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quasi-Fermi level of TiO2 and the electrochemical potential of the redox couple of electrolyte, mredox (qVOC ¼ EF - mredox); and FF the fill factor, a parameter (often evaluated by the “squareness” of the curve in Fig. 4.4C) which relates the Pmax (¼Jmax∙Vmax) to the VOC and JSC according to the relationship [45]: Jmax ,Vmax Pmax ¼ (4.3) JSC ,VOC JSC ,VOC The incident photon to current conversion efficiency (IPCE) (also known as external quantum efficiency (EQE)), defined as the number of generated electrons divided by the number of incident photons at a given l, is another significant performance metric. In more quantitative terms, its value corresponds to the JSC produced under (ideally) monochromatic illumination of the cell, divided by the incident monochromatic photons flux (Fl) that strikes the cell at that l. According to this definition, under short-circuit conditions, the IPCE (as a function of the incident l) can be determined as [45]:   JSC ðlÞ mA$cm2 IPCE ¼ 1240$ (4.4) lðnmÞ$fl ½mW $cm2  The IPCE value is also correlated to the light harvesting efficiency (LHE) (i.e., the fraction of incident photons absorbed by the dye) at the maximum wavelength l, to the injection efficiency of the electrons in the dye-excited state (Finj) and to the charge collection efficiency (Fcoll) of the injected electron at the photoanode according to [45]: FF ¼

IPCE ¼ LHEðlÞ$Finj $Fcoll

(4.5)

where the product, Finj∙Fcoll, represents the absorbed photons to current conversion efficiency, while LHE (¼ 1-10a, a being the dye’s optical density) depends on the dye’s extinction coefficient and concentration as well as on the absorbing layer thickness and the time of flight of photons within that film. Convolution of IPCE with the solar spectrum allows to obtain the generated photocurrent density in short circuit: Z JSC ¼ LHEðlÞ$Finj $Fcoll dl (4.6)

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It is worth to note that, IPCE can be considered as the collective contribution of the three fundamental mechanisms (Eq. 4.5), which must be synergistically optimized in order to raise the cell conversion efficiency. Good DSSC performances require IPCE value close to unit (approaching 100%), which implies that nearly every photon of selected l is harvested by the dye molecules (LHE w1) and that nearly every photoexcited dye molecule successfully injects an electron into the semiconductor nanoparticle network (Finj w1) and that electrons, diffusing through the semiconductor nanoparticle, are collected at the transparent electrode without recombining with the oxidized dye or adjacent oxidizing agent (Fcollw1). As a result, a suitable optimization of the interplay pathways of photoinduced electron transport and recombination mechanisms appears as the basic requirement of any artificial photosynthetic device. IPCE dramatically decreases with increasing the electron recombination rate (lifetime shortened). On the other hand, remembering that the charge injection is in furious competition with other radiative and nonradiative processes, the Vinj magnitude will be determined by the kinetic interlacement between these pathways [61]: Finj ¼

kinj kinj þ kdec þ kquen

(4.7)

where kinj and kdec are the rate constants for electron injection and the decay of the excited dye, respectively, while kquen include possible quenching routes of the dye due to aggregation, resulting in an enhancement of nonradiative decay [62] or reductive effects by oxidized redox couple on the excited dye state [7]. Also, the injection rate constant kinj, at solid/liquid interface into wide band gap semiconductors for an activationless process, can be derived from the Fermi’s golden rule [63]:
2 4p kinj ¼ (4.8) jHj2 rðEÞ h where h is the Planck’s constant, jHj is the electronic coupling constant between the wave functions of the donor (dye excited states) and the acceptor (CB of the nanostructured photoelectrode) with a value that is predicted to decline exponentially with the donor/acceptor distance because of the exponential radial character of their electronic wave functions and the CB’s density of states (r(E)).

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In DSSCs, the dye is strongly adsorbed on the semiconductor surface with the result of a large H between the dye’s excited p* orbital and the semiconductor’s CB, which possesses a rather large density of states. Thus, electron injection occurs at higher rates than the dye relaxation from the excited state toward its ground state, a feature that concurs to enhance the h. Since 2001, Asbury et al. [64] proved, from fs-IR spectroscopic measurements, an electron injection rate of w50 fs and 100 ps in TiO2 and ZnO thin films, respectively, both sensitized with an N3-ruthenium complex. The faster injection in TiO2 was rationalized in terms of higher density of states (DOS) and more favorable coupling with the electron-donating orbital (dye’s p* orbital). In fact, while the CB of ZnO possesses predominantly 4s-orbital character of Zn2þ, that of TiO2 comprises mainly unfilled 3d orbitals of Ti4þ, a feature which leads to a higher bulk density state (w190 times). Additionally, due to the different electronic structure of their CBs, the p* electron-donating orbital of the adsorbate overlaps more efficiently with the electron-accepting t2g-d orbital of Ti4þ favoring a more rapid electron injection. In 2002, Hara et al. [65] also explored the influence of dye’s molecular structure on the electron injection dynamics using two different Ru-phenantrolyne-based sensitizers, having four and one carboxyl groups, DCP2 and MCPP, respectively. The DCP2/TiO2 system showed a Vinj w 4 times larger than the other (MCPP/TiO2) where a single effective site for electron injection exists. In fact, by considering the electron in LUMO predominantly localized at the phenanthroline ligand, MCPP, differently from and DCP2, disposes only one site to anchor the TiO2 surface via carboxyl group, the other remaining as inactive remote ligand. Beyond other boundary contexts, such as the influence of electrolyte composition [66,67] or the sensitizer protonation state [61,68], several theoretical and experimental studies demonstrate the crucial role of both semiconductor and sensitizer nature in determining the electron injection dynamics [69e71].

4.3 Photoanode or working electrode According to the abovementioned working principles, it clearly appears that, for a DSSC to be successful, all the components must effectively play their role. Photoanode in particular should realize some critical requirements. First, the substrate for the semiconductor deposition should have a transparency 80% to permit the passage of optimum sunlight, together with a high electrical conductivity (s, (U/)) for efficient charge transfer

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and reduced energy loss in the DSSC. Transparent conductive oxide (TCO) on glass, indium-doped tin oxide (ITO, In2O3: Sn) with a transmittance (Tr) > 80% and a sheet resistance (Rsht) of w18 U/, and fluorine-doped tin oxide (FTO, SnO2: F), with a Rsht of w8.5 U/and Tr of w75%, are the wide diffuse substrates in DSSCs and many other optoelectronic applications. Second, the deposited semiconducting oxide should be a high porosity crystalline nanostructure, in order to offer high-specific surface area for dye loading, being also characterized by a wide energy band gap (3.0 O 3.2 eV). Although several wide energy band gap semiconducting materials (TiO2, ZnO, SnO2, NiO, etc.) and dispersant additives (e.g., organic polymers and carbonaceous materials) have been investigated as promising mesoporous metal oxide nanocrystal films [72,73], TiO2 nanocrystal, in the anatase allotropic form, remains the most popular one in DSSCs ambient, due to its best performance and its advantages in terms of costs, easy availability, nontoxicity, and chemically stability.

4.3.1 TiO2-based photoanode The choice of photoanode’s materials and their structure, determining the light harvesting, scattering ability, and charge transport, is among the main focus to engineering high-conversion efficiency DSSC. In this regard, photoanode based on a variety of nanosized crystalline materials in different geometries including nanoparticles (NPs) (zero-dimensional), nanowires (NWs), nanotubes (NTs) (one-dimensional, 1D), nanoflakes, and nanosheets (twodimensional structured materials, 2D) [74] have been extensively investigated and employed to boost the overall DSSC photoconversion efficiency, anatase-TiO2 porous phase, remaining yet the best performing choice. DSSCs with photoconversion efficiencies as high as 12% O 14% have been reached using TiO2-based photoanode [36,43]. Attention has also been paid to the TiO2 rutile form; however, the inability to preserve the nanocrystalline morphology at high synthesis temperatures poses potential issues for its exploitation in DSSC devices. Moreover, its lower energy band gap (w3.0 eV) compared to the anatase (w3.2 eV) makes this latter chemically more stable [75]. Nanostructured ZnO, with a band structure similar to TiO2 and higher electron mobility compared to the latter, together with a myriad of options to synthesize the more appropriate shape and size, has been considered as a promising photoanode materials. However, despite its favorable peculiarities, ZnO exhibits higher recombination rate, lower dye loading, and

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slower electron injection compared to TiO2 [76]. Moreover, in acidic environment, ZnO shows instability and formation of dye aggregates on its surface causing a downfall in the cell performances [77], probably as a result of its acidic surface and more positive CB edge (3.37 eV) compared to the TiO2. SnO2 has also been considered as a further attractive option in photoanode materials, due to its carrier mobility (w200 cm2V1s1 [78]) higher than that of TiO2 (w0.1 O 1 cm2V1s1 [79]) together with a larger band gap (w3.8 eV) which should be guarantee of fewer oxidative hþ generation in the VB under UV illumination (lower dye degradation rate and better long-time cell stability) than those made from TiO2 or ZnO, facilitating also the electron injection from photoexcited dye [80] (more positive band edge position than TiO2). However, the PV performances of DSSCs exploiting SnO2-based photoanode and Ru-complex dyes (N719) appear surprisingly lower than that based on TiO2, a feature due to the concurrence of several adverse factors such as unfavorable band alignment between the two partners and faster electron recombination dynamics [81], as well as a lower dye adsorption degree [82]. Although strategies have been developed to overcome the above issues, such as treating of the SnO2 photoanode with TiCl4 prior to anneal (thus inducing a TiO2 layer atop the SnO2) or the use of a mixed photoanode (SnO2/ZnO [83]), these approaches have so far failed to overtake the best TiO2-based devices performance. Other semiconductor materials, such as WO3 [84,85] and Nb2O5 [86] were also studied. Fig. 4.5 shows the band gap values of some common semiconductors versus vacuum and NHE energy reference (NHE at 4.5 eV with respect to the vacuum level). Following the innovative concept by O’Regan and Gr€atzel [27] in the use a mesoporous film of small NPs as scaffold for light harvesting in PV devices, TiO2 photoanode, has been immensely studied and, various approaches and photoanode architecture have been developed to optimize its functionality (i.e., reduced charge recombination, electron transport, dye pickup, and light scattering ability). A double layer architecture, where an additional layer of relatively large size overlaps a TiO2 nanostructured underlayer, has been suggested to improve the layer-scattering and consequently the photoanode’s absorption/scattering ability. Mohammadi et al. [87] reported a double-layer DSSC (TiO2 NWs/TiO2 NPs) with reliable conversion efficiency of 6.34%, arising from a fast electron transport and lifetime, together with low photoelectrons recombination and efficient dye adsorption and light scattering.

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Figure 4.5 Band gap energy, VB and CB positions of several common semiconductors on a potential scale (V) versus NHE/vacuum.

In 2007, O’Regan et al. [88] proposed that the treatment of TiO2-based photoanode with TiCl4 (heating to 450  C) produces large reductions in the electron/electrolyte recombination rate constant (up to 20-fold) improves the binding of the dye, increasing injection efficiency, and also shows some possible influence on the charge transport. The combination of all these factors enhances the overall DSSC h. The decrease in the recombination rate constant could also be correlated to downward shift in the TiO2CB with an increase in charge separation efficiency for absorbed photons and a consequent photocurrent rise. The use of 1D nanostructured materials (NWs and NTs) has been proposed as a profitable straightforward path to electron transport, able to minimize the time inside the photoanode (recombination probability) while maximizing the charge collection, this latter directly connected to the electron diffusion length Ln: at longer diffusion length corresponds higher charge collection efficiency. Unfortunately, NW-based cells show low dye loading and a consequent low optical density, which limits the photocurrent density. Employing composite materials showing both high-specific surface (e.g., mesoporous TiO2-NPs) and fast electron transport (e.g., singlecrystalline TiO2-NWs) has been suggested as a solution to overcome this drawback [89,90]. TiO2 nanosheets (NSc) have also been applied to obtain good electron collection together with high-specific surface and reliable light scattering ability when composite with TiO2 NPs. h of w10%,

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significantly higher compared to standard P25-based photoanodes (7.8%), has been reported for optimized photoanodes [91,92]. Once again, it clearly appears that the crystalline and electronic structures (related to the chemical nature of the bonds between atoms or ions, atomic arrangement and physical dimensionality) of TiO2 materials, determining largely their photoactivity, remain essential issues for their improvement and exploitation in the current and future energy nanotechnology. Structurally, TiO2 occurs in three polymorphs in nature, namely rutile (tetragonal D14 4h -P42 mnm, a ¼ b ¼ 0.4584 nm, c ¼ 0.2953nn), anatase 19 (tetragonal D4h -I41 amd, a ¼ b ¼ 0.3733 nm, c ¼ 0.937nn), and brookite (rhombohedral D15 2h -Pbca, a ¼ 0.5436 nm, b ¼ 0.9166 nm, c ¼ 0.2953 nn) [93] in order of decreasing abundance. Rutile and anatase being widely investigated and extensively characterized for their role in TiO2 applications. Their structures are based on a tetragonal Bravais lattice, simple in the rutile, body centered in the anatase, each Ti4þ cation always surrounded by an octahedron of six oxygen anions (O2) and each anion is coordinated to three cations leading to an OTi3 pattern [94]. The resulting structures can be described as chains of distorted TiO6 octahedra, with two long “axial” and four short “equatorial” TieO bonds. The main differences being the octahedral distortion degree and in the way how moieties are assembled to build crystal (e.g., angles and bond lengths (see Fig. 4.6A)). The octahedra in anatase is more distorted than in rutile, with a generous deviation from a 90 degrees bond angle between the titanium and oxygen atoms in anatase. In addition, in rutile, each octahedron shares corners (in the plane 001) with eight neighbors, and shares edges with two other neighbors (equatorial oxygen atoms), so each octahedron is in contact with 10 neighbor octahedrons in a linear chain disposition. On the contrary, in anatase each octahedron shares corners with four neighbors and shares edges with four other neighbors, so each octahedron is in contact with eight neighbors in a zigzag chain structure with a screw axis [95,96]. These differences, beyond the different mass densities (4.13 g/cm3 and 3.79 g/cm3, for rutile and anatase, respectively), give rise to different electronic band structures which, in turn, affect the material’s optical response. Extensive studies have shown the existence of severe distortions in the Ti site environments and more effective photoactivity, when the size of TiO2 particles becomes smaller than 20 nm [96e98]. In these circumstances, the surface Ti atoms change their coordination environments from octahedral to pentacoordinate square-pyramidal (see Fig. 4.6B) with the consequent TieO bond length distortion, depending on the nanoparticles curvature.

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Figure 4.6 (A) Lattice structure of TiO2 and arrangement of TiO6 polyhedral in anatase and rutile polymorphs. The axial and equatorial bond lengths and angles of the octahedrally Ti atoms are also indicate. Large red spheres represent the oxygen atoms, smaller gray spheres indicate the Ti atoms. (B) Schematic representation of the titanium coordination environment in flat TiO2 surface and small TiO2 particles. (c) Total density of the states of anatase TiO2 (top) and projected DOS for Ti 3d and O 2p orbitals (below). The top of the VB (solid line) is taken as the zero of energy. The dashed line indicates

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In such a lower coordination structural degree, extremely small TiO2 particles expose at their surface a large number of under-coordinated “corner defects” which are the source of novel enhanced and selective reactivity of the TiO2 NPs toward surface-attached ligands, inducing new hybrid features and significant alterations in the electrochemical properties of the modified NPs. Ligand attachment of organic compounds at the TiO2 surface (chemisorption) may be considered as an acid-base reaction where a donor group containing oxygen, nitrogen, or sulfur atoms, acts as a Lewis base and donates electrons to the Lewis acid (surface TiIV). On the other hand, this under-coordinated “corner defects” (oxygen deficiency) can also be considered as Ti3þ impurities providing n-type doping characteristics to the TiO2 semiconductor. On the other hand, the materials optical response will also depend on their underlying electronic structure, i.e., on their DOS closely connected to the nature and length of chemical bonds as well as to the size of the structural arrangement (confinement of carriers). The DOS of TiO2 consists of Ti eg, Ti t2g, O ps, and O pp (Fig. 4.6C), the oxygen 2p states bounded to titanium’s 3d-orbitals being the main contribution to the VB of both bulk TiO2 and its nanostructures, while the CB comprises the titanium 3d eg states (5 O 8 eV) and t2g states (2O5eV) involved in an antibonding interaction with oxygen 2p orbitals [99], with the dxy states of t2g bands predominantly located at the bottom of the CB (dashed line in Fig. 4.6C). Also, it is worth to note that the lower part of TiO2 CB, consisting of d orbitals of t2g symmetry, is expected to allow strong electronic coupling with electron donating groups [100]. On the other side, the TiO2 upper VB can be decomposed into three main regions: (1) the s bonding in the lower energy region where oxygen ps mainly contribute to the bonding, (2) the p-bonding in the middle energy region, weaker than that of the s-bonding, and (3) oxygen pp states in the higher energy region where the hybridization

=the CB minimum as a guide to the eye. (D) Molecular orbital energy level diagram of anatase TiO2: the occupied states of VB and those of the CB edges arise predominantly from the O 2p and Ti 3d orbitals, respectively. The octahedral coordination causes a crystal field splitting of the 3d states into the t2g (dxy, dxz, and dyz) and eg (dz2 and dx2-y2) subbands. The eg orbitals point toward the O atoms forming s-type orbitals. The t2g point is in between the oxygen neighbors and form p-type bonds. (B) Adapted from U. Diebold, The surface science of titanium dioxide, Surf. Sci. Rep. 48 (2003) 53e229; S.M. Gupta, M. Tripathi, A review of TiO2 nanoparticles, Chin. Sci. Bull. 56 (2011) €tzel, Photochemical H2 generation by 1639e1657. (D) Adapted from V.H. Houlding, M. Gra visible light. Sensitization of TiO2 particles by surface complexation with 8-hydroxyquinolin, J. Am. Chem. Soc. 105 (1983) 5695e5696.

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with d states is almost negligible, leaving the oxygen pp state nonbonding at the top of VB. High-quality theoretical and experimental studies have been performed to understand the relationship between structural morphology and interband electronic transitions in order to better design TiO2 photoactive materials. Already in 1999, Braginsky and Shklover [101] suggested for small porous semiconductor crystals (few 10 nm or less in size), where the share of the interface atoms is sufficiently large, the possibility of momentum nonconservation at the interface and a considerable enhancement of the light absorption. A favorable interfacial mechanism with enhanced absorption in indirect semiconductors (e.g., TiO2) where the direct electron transitions between the band centers are prohibited by the crystal symmetry. Later, Sasaki [102], Sato [103], and Sakai [104] investigating TiO2 on ultrathin nanosize dimensions (nanosheets) showed that in these cases the band gap was larger than that of anatase TiO2, as a consequence of a lower dimensional morphology of the former (3D / 2D transition). Furthermore, on the basis of their results, Sakai and coworkers determined values of 3.8 eV (equivalent to an absorption edge of lg z 330 nm) and 3.2 eV (corresponding to an absorption edge of lg z 390 nm) for the band gap energy of TiO2 nanosheet and anatase, respectively, and flat-band potentials positioned at 1.27  0.05 V and 1.15 V  0.05 V. In practice, the titania nanosheets absorption band appeared significantly blueshifted with respect to bulk TiO2 anatase. Unlike other wide band gap semiconductors of similar band gap and band edge potentials such, as ZnO and SnO2, whose CBs are constituted of empty s and p orbitals of Zn2þ and Sn4þ ions, the TiO2 CB is comprised mainly from empty d orbitals of Ti4þ ions, with d bands which are typically narrower and have DOS orders of magnitude higher than sp bands [105]. Thanks to this intriguing electronic structure, TiO2 remains an attractive prototype material on which the fundamental issues of semiconductor electrochemistry can be developed. In addition, with appropriate ligands, titania may form complexes with strong electronic coupling. Fig. 4.6D depicts a molecular orbital bonding diagram, derived from the discussed character of the states. Despite its many applications, the wide band gap of TiO2 polymorphs (3.0 and 3.2 eV for rutile and anatase phases, respectively), covering a small fraction of the sun’s energy ( carboxylic acid > ester > acid chloride > carboxylate salts > amide. Sensitizers with D-p-A architecture, based on a diphenyl/dialkylamino electron donor group (D) and carboxyl group as electron acceptor (A), are promising candidates for type-2 DSSCs due to their easiness to admit suitable substituents onto the chromophore skeleton, allowing to adjust the dye’s stereochemical structure as well as its photophysical/electrochemical properties (HOMO and LUMO energy levels tuning). Remarkably, in D-p-A dyes the HOMO is localized over the p-conjugated system close to the D moiety, and the LUMO is localized over the A moiety [114,115]. Also, D-p-A dyes facilitates the intramolecular charge transfer from the D to the A moiety and give rise to both efficient electron transfer (from the excited dye to the TiO2 CB) and D-A charge separation in the excited state that implies efficient electron transfer and retarded charge recombination. Due to these benefits, much effort has been devoted to advance D-p-A dye sensitizers for type-1 DSSCs [116e118].

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On the other hand, numerous natural pigments, such as bromopyrogallol red and anthocyanins containing catechol (Cth, C6H10(OH)2) moiety, exhibit a direct electron injection pathway from the dye’s HOMO level to the CB of the TiO2 by photoexcitation of the dye-to-TiO2-charge transfer bands (DTCT). The major advantage of this pathway is its lightharvesting ability over a wide sunlight spectral range, since in this case the direct electron injection can lead to generate a broad DTCT band and to attenuate the restrictions on the LUMO showed by an indirect electron injection route (type-1 DSSCs). However, though Cth-TiO2 electrontransfer dynamics have been widely investigated [119,120], the PV performance of type-2 DSSCs is about 10-fold lower than that of the based on “two-step electron injection” pathway [121]. One key of this limiting factor could lie in back electron-transfer dynamics (charge recombination) from the reduced TiO2 to the oxidized sensitizers, which is faster in the onestep route than in the type-2 DSSCs [122]. However, in 2005, Tae et al. [123] proved that the presence of electron-donating groups on the Cth moiety can retard the back electron-transfer kinetics as well as can increase the electron injection efficiency, causing also a redshift of the DTCT band, due to a higher donor strength of the dye and the consequent increase in the visible light absorption. Their results revealed, for the first time, the potentiality of efficient Type-1 DSSCs. Cht has been known to bind to the TiO2 surface either by dissociative absorption (coordination) or by molecular absorption, the latter involving formation of van der Waals complex via hydrogen bonds. In dissociative absorption, three different structure may be adopted: (1) the monodentate, where only one of the oxygen atoms is bonded to a titanium atom on the surface, (2) a chelating bidentate structure, where both oxygen atoms are bonded to the same titanium atom on the surface, and (3) a bridging bidentate structure, where each oxygen atom is bonded to a different titanium atom on the surface. Theoretical studies by Redfern et al. [124,125] revealed that when Cht adsorbs on a stoichiometric TiO2 surface, its bridging bidentate structure is favored by 4 O 6 kcal compared to the molecular absorption, while the chelating bidentate structure prevails when absorption occurs at a defect site (containing a Ti]O moiety). Photoemission results suggest that the dopamine molecule adsorbs on the surface in a bidentate geometry (i.e., bonding through both oxygen molecules). In 2009, Li et al. [126], by combining UV-photoemission spectroscopy (UPS), scanning tunneling microscopy (STM) and DFT calculations, proved that Cth forms ordered

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densely packed monolayer on the TiO2 surface giving rise to electronic states in the band gap above the TiO2 VB edge. Also, by DFT analysis, the presence of either monodentate only or mixed monodentatebidentate structures, which can easily convert from one into the other via proton exchange between the surface and the adsorbed Cth, was also asserted. In 2010, Thomas et al. [127] by photoemission data and computer modeling showed that dopamine (a correlated catechol molecule) binds two surface titanium sites in a bidentate tilted geometry with the phenyl rings orientated normal to the TiO2 surface. In 2012, Kaniyankandy et al. [128], investigating the influence of the ligand molecular geometry on the interfacial electron transfer dynamics in enediol-TiO2 systems (catechol, resorcinol, and quinol), observed changes in enediols/TiO2 electron coupling depending on the eOH binding groups location. A larger eOH distance (resorcinol and quinol) causes the two Ti4þ ions, involved in a bidentate binuclear binding arrangement, to be more distant each to other, consequently electrons will be delocalized over a larger area and lower recombination dynamics will be expected (see Fig. 4.7).

4.4 Natural dyes An ideal sensitizer should broadly absorb across the visible and near-IR portion of solar spectrum with high molar absorptivity. Simultaneously, it should show fast injection of the photogenerated electron from its excited state to the proximal semiconductor’s CB (when compared to the lifetime of any other relaxation or quenching processes) [45]. Also, it must bind strongly to the semiconductor oxide surface, as well as possess a suitable

Figure 4.7 (A) Schematic depiction of “one-step” (left) and “two-step” (right) electron injection pathways from an adsorbate dye to the TiO2 surface. (B) Possible binding modes of catechol group on TiO2 surface and (C) enediol structures.

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redox potential for rapid regeneration through the redox mediator and, show high photostability (capability to sustain many reduction/oxidation cycles without decomposition) under illumination condition and continuous light soaking. Ideally, for DSSC applications the dye should not show notable degradation effects for an operation time as long as 20 years [129]. However, though electron injection and regeneration can be considered as reversible pathways, in a real device, some degradation degree of the dye’s takes place always, with a rate constant (k) according to the following [129,130]. k ¼  ln

Cdye;t scell Cdye;0

where k (s1) is the degradation rate constant, Cdye,0 and Cdye,t are the dye concentration at the time zero (initial) and “t,” respectively, and scell the cell lifetime (20 years z6.3$108 s). In addition, since injection is a very fast process, it is plausible to assume that light-induced degradation occurs mainly from the dye oxidized state (Dþ) rather than from its excited state (D*). Efforts have been made to develop efficient dye sensitizers, broadly classified in three main families: (1) metal complex dyes, (2) metal free organic dyes [131], and (3) natural dyes. Structurally, metal complex dyes have two ligands, anchoring and ancillary, the first being responsible for absorption on the semiconductor surface, whereas the ancillary is determinant to tune the complex’s overall properties [132]. Ruthenium bipyridyl derivatives (N3, N719, and N749 or black dye) have achieved the high conversion efficiency in this class of sensitizers. However, the high costs, intricate synthesis route, and limited extinction coefficient (usually < 2.50∙104 M1 cm1 in the visible region) have restricted the developing potential of Ru-dye-based DSSCs. Metal-free organic dyes have also show great promise with good flexibility in molecular design, stability under prolongate light condition, and extremely high molar extinction coefficient, at least one order of magnitude higher than Ru-complexes, making the thin TiO2 layer feasible. Structurally, metal-free organic dyes are donor-p-acceptor (D-p-A) configuration consisting of an electron donor and an electron acceptor, which are linked covalently through a p-conjugated bridge/spacer. Usually, coumarins, phenylamines, indolines, triarylamines, and carbazoles are used as electron-rich moieties, while cyano-acrylic acid, rhodamines, and pyridines are employed

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as acceptor part. The linking between donor and acceptor is selected from thiophene, polyene, and benzothiadiazole derivates [133,134]. Natural dyes are mainly derived from vegetable sources (e.g., fruits, flowers, leaves, and so on [45]). The use of natural pigments (e.g., chlorophyll, betalains, anthocyanins, carotenoids, etc.) as sensitizing dyes in DSSCs ambient, has come to be a worthwhile substitute to the inorganic/organic sensitizers because it enhances the cost-effective and, more importantly, does not pose environmental pollution risks. In addition, natural dyes show a broad colors range and strong absorption in the blue and red regions of the visible spectrum together with excellent redox properties and anchoring functional groups that interact favorably with the surface of TiO2 semiconductor. Up to now several natural dyes have been used in DSSC with appreciable level of PV performances [45,135e137]. In the following, we present an in-depth analysis of the most significant natural pigments for future DSSCs development.

4.4.1 Chlorophylls Chlorophylls (Chls) are the key components of natural photosynthetic systems in green plants, bacteria, and algae. In this respect, Chls and their analogues, whose spectral and photochemical properties were widely investigated [138,139], have been explored as natural sensitizers in DSSCs, thank to their efficient light-harvesting mechanism and electron transfer reactions. In 1971, Tributsch and Calvin [140] first investigated the photosensitization of a wide-band gap ZnO semiconductor by Chls pigments, and later other research groups [141,142] exploited Chls analogues as sensitizers to extend the photoresponse of colloidal TiO2 semiconductor. Chls, refers to a class of cyclic tetrapyrrole, likewise to porphyrin, with a magnesium ion at center [143] and, in this regard, porphyrin macrocycle can be considered the core of Chls structure (see Fig. 4.8). Porphyrins have been studied intensively by computational and spectroscopic measurements. Ab initio calculations (restricted Hartree Fock, local density functional theory) have shown that both free base and metal (Mg(II), Cu(II), Zn(II), Ni(II)) containing porphyrins generally possess two a1u- and a2u-type orbitals as the HOMO and HOMO-1 [144] with the electron density in a1u orbital essentially localized on the Ca and Cb atoms of the pyrrole rings and, in the a2u orbital mainly found on the pyrrole nitrogen and meso carbon atoms [145]. Likewise, HOMO, the LUMO, in both metal-free and metalporphyrins, show two nearly degenerate eg(p*) orbitals essentially centered

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on the porphyrin macrocycle. Also, the frontier orbitals justify the similarity in absorption spectra of porphyrin derivatives. Chls and homologous compounds show two characteristic pep* absorption region: the Soret-band located in the blue region (400 O 450 nm), deriving from the electronic transition with the strongest oscillator strength (2.88 eV) and the Q-band lying in the red region (550 O 700 nm), consisting of Qx and Qy transitions (see Fig. 4.8) at 1.87 and 2.14 eV, respectively [146]. In addition, Chls displays ε as high as 105 M1cm1 and capability to tune both their Q-bands position/intensity and excited state lifetime via macrocycle chemical modification and by changing the central metal, respectively [146]. Generally, Chl-a derivatives (named also as chlorophyllins) display greater sensitizer ability than pure untreated Chl-a [142]. A feature explained by considering the phytyl-ester group interferences, which prevent effective bond between the carboxylic groups of the Chl-a and the TiO2 NPs. In this regard, numerous studies on chlorophyllin as sensitizer into solar energy applications have been focused on Chlorin-e6 (Chl-e6), obtained by the alkaline hydrolysis of Chl-a, leading to the opening of cyclopentanone ring and the creation of two additional carboxylic groups [142]. In 2007, Amao and Yamada [147] proposed a Chl-a derivative zinc chlorine-e6 (ZnChl-e6), assembled in hydrophobic domain onto

Figure 4.8 Absorption spectra of chlorophylls (A and B) together with their molecular structures. A schematic description of pep* transitions (responsible for the Soret and Q bands) are also shown.

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nanocrystalline TiO2 electrode. The use of a co-adsorbed cholic acid strongly suppressed the dye molecule aggregation by generation of large hydrophobic domains of rigid skeleton of cholic acid, and improvements in the DSSC performance compared to the cell using merely ZnChl-e6 based electrode were reached. In 2010, Chang et al. [148] used spinach and ipomoea extracts, as natural sensitizers for DSSC photoelectrodes. However, the h of the DSSCs based on these kinds of sensitizers were rather low due to the absence of effective anchoring groups, such as carboxylic group, that directly conjugated to p-electrons of the chromophore, allows efficacy electron injection into the TiO2 and, in this respect, Chl-c better meets these criteria. Chl-c has a carboxylic terminal group to anchor the TiO2 NP having also a structural arrangement responsible for a large electron density conjugated with the carboxylic group and consequently, capability to inject electron from the porphyrin macrocycle to CB [143]. The order of conversion efficiencies in TiO2-based solar cells employing different types of Chl-c (c1 and c2) and their oxidized forms (c1’ and c2’) isolated from a brown seaweed (Undaria pinnatifida) determined in the order, Chl-c2 > Chl-c1 > Chl-c2’ w Chl-c1’, has been explained in terms of the more extended conjugated chain and the resultant higher electron densities in Chl-c2 than in Chl-c1, as well as of the different configurations of the vinyl carboxyl group to form a hydrogen bonding-like p-electrondHO interaction in Chl-c2’ and Chl-c1’ [149]. Among the cited Chls, higher absorption on TiO2 surface and larger IPCF (w80% at 425 nm) were showed by Chl-c2. In addition, the absorption spectra of Chl-c2-based photoanode appeared redshifted compared to the corresponding electronic spectra in solution, as a result of macrocycles polarization due to an electron flow toward TiO2dinstead that a dye aggregation through a stacking of the porphyrin macrocycles. Strong coupling between the sensitizer’s delocalized p electrons and the nanostructured semiconductor together with suitable energy level of the dye HOMO level and broad/intense Qy band, as well as presence of carboxylic groups, alkyl chains, and electron donors, are key requirements to enhance the overall solar energy to electrical power conversion with these kinds of sensitizers. In this framework, the p-conjugation length extension along the y-axis could be an exploitable strategy to increase the dye extinction molar coefficient and to shift the HOMO orbital to a higher energy level. Location of eCOO groups along the y- or x-axis direction (Fig. 4.8) gives rise to strong coupling chromophore/TiO2. Differences in y-/x-directions influence the absorption properties and the molecular orbitals of Chl sensitizers.

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Light-harvesting ability of Chl sensitizer to improve the DSSCs performance has been explored by comparing different Chls sensitizers having electron-withdrawing p-conjugatable eCCO groups at different positions on different molecular axis of the Chl macrocycle (see Fig. 4.8). Much denser Qy band in Chl macrocycle with withdrawing p-conjugatable eCCO group at the C3-position (Chl-1 in Fig. 4.9) on the y-axis, showed much denser Qy band than the counterpart having eCCO group at C3- or C7- position on the x-axis (Chl-2 and Chl-3 in Fig. 4.9). Thus, DSSC based on Chl-1 gives higher photocurrent and performance than those based on Chl-2 and Chl-3 [143]. In 2014, Calogero et al. [150] performed an interesting study on Chl-e6and copper Chl-e6 (Cu Chl-e6)-based DSSCs. The influence of metallic ion and solvent polarity on the dye aggregation and the electrodes absorption bands were analyzed by electronic absorption measurements and compared with previous literature data. Red and blueshifts of the Soret- and Qy-peak, respectively, in the simple Chl-e6-based electrode were ascribed to a higher solvent polaritydwhich influencing mainly the dye LUMO level, reduces the HOMO-LUMO energy gapdand to the dye aggregation process, respectively. Differently from simple Chl-e6, in copper Chl-e6 no significant solvent effect appeared for both Soret- and Qy-peaks. A circumstance explained by considering a higher porphyrin ring electronic delocalization, due to the Cu d orbitals, which makes copper-Chl less polarizable than the counterpart Chl-e6. On the other hand, the Qy wavelength appeared strongly affected by the specific central metal, showing a blueshift which increases with increasing the central metal electronegativity. A feature in agreement with the evidence of previous analysis about the existence of electrostatic interactions between the central metal and the N-centered

Figure 4.9 Molecular structure of Chl-1, Chl-2, and Chl-3 with carboxyl groups located at different positions on the Chl macrocycle.

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p-electron densities is the capability to modify the energy levels of the electronic transitions of the system. The better light harvesting ability in the Qy wavelength region together with a higher IPCE, in both Soret- and Qypeaks, showed by the copper Chl-e6 photoanode compared to the Chl-e6, allowed the authors to propose as a rule that with increasing the Soret-Qy wavelength difference, a decrease in DSSC efficiency will be expected.

4.4.2 Betalains Betalains are water-soluble nitrogen-containing pigments, early erroneously addressed as flavocyanins and nitrogen anthocyanins. The current name was coined, from the Latin (Beta vulgaris) by Mabry and Dreiding [151]. Beyond their nutritional perspectives [152], betalains pigmentsdoccurring prevalently in caryophyllales and in some basidiomycota plants [153]drepresent an additional class of potential dyes due to their high molar extinction coefficients in the visible region, good light harvesting capability, pHdependent redox properties, and facility to coordinate metal ions. To date the betalains comprise a quite modest number of about 55 structures [154,155]. From a chemical point of view, betalains derive from betalamic acid, and depending on the components bonded to the main structure, betacyanins or betaxanthins arise, the former when the group is 3,4 dihydroxyphenylalanine (DOPA), which may or may not be glycosylated, and the latter if the conjugation partners are amino acids or derived amines (see Fig. 4.10). The red-purple betacyanins (betanin (I) and betanidin (II)) show a maximum absorptivity at lmax of w535 nm while the yellow betaxanthins (vulgaxanthins I and II and indicaxanthin III) have a lmax around 480 nm [156,157]. Aglycone betanidin shows a slightly redshifted absorption spectrum than betanin. Table 4.1 lists the lmax values and molar absorptivity ε for common beet pigments [158].

Figure 4.10 Molecular structure of betalamic acid and betalain derivates.

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Table 4.1 - Wavelengths of maximum absorption, lmax, and molar absorptivity, ε, of common beet pigments. Dye l (nm) ε (ML1cmL1)

Betanin Betadin Indicaxanthin Betalamic acid

535 542 482 424

65,000 54,000 42,600 24,000

Given their strong visible electronic absorption, disposition toward electron donation, and existence of anchoring eCOOH functional group in their main structure, betalains could be promising dye sensitizers to exploit as excited state electron donor in TiO2-based DSSCs. At the best of our knowledge, there are few reports published about the use of betalains-based DSSCs, the main research being mainly focused on chlorophyll and anthocyanins. Here we reviewed some results of an interesting experimental work on raw extracts of red beet, Sicilian prickly pear fruits, and bougainvillea. The red beet (i.e., red turnip or beetroot) possesses ball-shaped red roots containing high concentrations of betalains (order of g/kg) [159], while prickly pear contains both betanin and indicaxanthin. Finally, the color of bougainvillea plants is due to the combination of the two main dyes: betacyanin and betaxanthin. In 2002, Zhang et al. [160] assessed, for the first time, betalain pigments (from beet red, purity of w80%), stabilized with sodium L-ascorbate, sodium metaphosphate and dextrin, as dye sensitizers for TiO2- and ZnO-based DSSCs and, preliminary hs as low as 0.44% and 0.014%, respectively, were obtained. The lower performance of ZnO-based electrodes was attributed to a less electron injection efficiency from excited states to the ZnO CB and/or a higher recombination of the injected electron than TiO2-based electrodes. Betalains, via their carboxylic group, strongly bind the TiO2 surfaces forming an ester-type linkage [161]. The p*-orbital of the carboxylic group promotes a faster excited-state electron injection into the CB of TiO2 than in that of ZnO due to the different symmetry. For TiO2, it is conceivable that the overlap between the extended p-system and titanium 3d orbital can aid the electron injection more efficiently than the mixing between carboxylate p-system and the s-orbital in the ZnO [162]. In 2007, Quin and Clark [163] performing density functional theory (DFT) and time-dependent functional theory (TDDFT) calculations, in betalain pigments, demonstrated that the HOMO is mainly localized on the molecule’s benzene ring. As a result of the optical transition, a loss in

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the electron density and a charge gain in the benzene and dihydropyridine (DHP) rings, respectively, have been proved, indicative of partial excited state electron transfer pep* transition. By the assumption that the DHP unit is anchored to the TiO2 surface via carboxylic groups, the excited state will necessarily be electronically coupled with the acceptor states of TiO2, generating an effective charge injection. According to Quin results, the betacyanin’s relative excited state oxidation potential vary in the 1.20 e 1.33 V (vs. NHE) range, consistently more negative than the TiO2’s CB edge (- 0.5 V vs. NHE), thus ensuring sufficient driving force for the photoinduced electron transfer. Similarly, the calculated betanin’s ground state oxidation potential (1.0 V vs. NHE) parallels the differential pulse voltammetry measurements [164], confirming that the oxidized betanin can easily be reduced by the iodide mediator ensuring good in dye cation regeneration. In 2008, Zhang and coworkers [158] reported the use of different betalains (e.g., betanidin, betaxanthin, and melanin), extracted from the red beet, as sensitizers for DSSC, betaxanthin showing the best sensitization activity. Interestingly, despite its low light harvesting capability, melanin-based DSSC produced a good current density, while betanin-based device exhibited the poorest efficiency depending on a low current density. An explanation to this behavior could be the presence of the cyclo-DOPA and glucosidic group, which gives rise to a large steric hindrance in the betanin molecule compared to the betaxanthin, a feature that is expected to influence the dye molecules adsorbed onto the anode. Also, the condensed ring in betanin molecule, delocalizing the positive charge could be responsible of the produce lower current density. Also, it was demonstrated that acid pretreatment of anode permitted higher dye loading and contemporary reduction of recombination of TiO2 electrons. A feature which clearly appears by comparing the absorption spectra of both natural pigments reported in Fig. 4.11AeB [145]. The high absorption below 450 nm arises by the betalamic acid, while above this spectral region the shape of the absorption curve reveals the same betanin in both extracts and different betaxanthins: vulgaxanthin I in red turnip and indicaxanthin in prickly pear. Red turnip absorption spectra appeared blueshifted with respect to prickly pear, a behavior also paralleled by the photoaction spectra (see Fig. 4.11C). The IPCE maximum of both red turnip and prickly pear were of w65% in the 450 e 470 nm region. By considering the negative excited state oxidation potential around 1.3 V together with good electronic coupling with the d band of the TiO2 promote “activationless” injection into the TiO2

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Figure 4.11 (A) Absorption spectrum of Beta vulgaris Rubra (red turnip) and (B) Opuntia engelmannii (prickly pear)-based TiO2 photoanode acidic treated. (C) Incident photon to current conversion efficiency curve of red turnip and prickly pear.

CB, the charge injection is not the limiting IPCE process. On the other hand, recombination losses can reduce the 4collection, while, the betalain cation reduction is expected to be faster than ruthenium artificial dye [145]. Concerning the stability of betalains by Sicilian prickly pear, preliminary test on sealed cells of 1 cm2 reported efficiency constancy for a time as long as 20h under one sun irradiation [165]. The thermal stability was also evaluated by calorimetric analysis, and decomposition processes above 160 C were observed. Finally, the presence of ascorbic acid has been proved to preserve betalains from thermal degradation. In 2013, Hernandez-Martinez et al. [166] increased the stability of modified betalains extracts from Beta vulgaris (red beet) over 72 h of working operation. In 2017, Wendel et al. [167] performing betanin-, vulgaxanthin I-, and others betalains based-DSSC reported insufficient electron injection quantum yield as mainly responsible of their low h with respect to the artificial ruthenium dyes. In fact, in betalains based-DSSC, just w50% of the excited dye molecules inject electrons, the other suffering fast internal conversion process from the S1 state. This latter route is enhanced by intramolecular

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rotation process (betalain acts in nature as good heat-release systems), which competes with electron injection pathway. Differently, despite in betaxanthin-based DSSC electron injection quantum yield appears higher (about 70%); however, the total sunlight conversion efficiency is smaller due to the blueshifted absorption band. A probable explanation of this difference in the electron injection quantum yields is due to the different kinetics of the competing S1/S0 internal conversion processes. Generally, dyes with high band gap and slow internal conversion rate display high electron injection quantum yield but reduction of long-wavelength absorption, this latter resulting in a gradual decrease of sunlight conversion efficiencies. Despite their high absorption coefficients and proper anchoring groups, betalains seem to be not suitable for DSSC applications due to their short S1 state lifetime. Engineering of more appropriate molecular architectures capable to hinder the molecular rotations in such kind of sensitizers should be undertaken.

4.4.3 Anthocyanins 4.4.3.1 Chemistry and equilibrium structures Anthocyanins (Ant) belong to a widespread class of phenolic compounds collectively named flavonoids characterized by a noticeable interaction with solar light. Due to their engaging optical activitydresponsible for most of the color diversity found in plants and flowersdchelating capability toward metal ions and visible light-harvesting capability Ants have been exploited as sensitizer in optoelectronic nanodevices including DSSCs [45,168e171]. As a class of compounds, anthocyanins have a diverse range of molecular structures that play important roles in determining their molecular electron distribution, reactivity, and color variety [172,173]. Differences in the number and position of hydroxyl and/or methoxy groups as well as the type, identity, and position of sugars attached to the carbon skeleton, which are able to be further acylated (acyl glicoside) with one or more acid moieties, determine the existence of the multitude of naturally occurring anthocyanins [45,173]. In addition, depending on environmental conditions, such as pH, temperature, light excitation, oxygen, metal ions, Ants undergo structural modifications accompanied by changes in the visible absorption spectrum as well as in the chemical stability. In aqueous ambient, between pH 1e8, Ants exhibit a rather complex equilibria between two colored species, flavylium cation (AHþ) and quinoidal base (Q), and two colorless ones, hemiketal (H) and chalcone (Ch) that can be driven by pH change. At pH  3, anthocyanins exist predominantly

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in the stable reddish colored flavylium cation form AHþ which is a weak acid (Ka ¼ 104O105) [174]. With increasing pH value (3 < pH < 6), most acid groups of AHþ (C40 , C5, and C7) undergo a rapid proton loss (preferentially at the eOH group at the C40 position) leading the red color toward a blue mixture of neutral tautomeric Q-base. A further pH-basic value increase (pH > 8) can induce a second deprotonation (usually in position C7) and the formation of a negatively ionized base, Q-, with the appearance of possible bathochromic and hyperchromic effects in the absorption spectra. Beside fast proton-transfer equilibria, AHþ also undergoes water addition at the C2 position with formation of a colorless H pseudo-base, which in turn can equilibrate with minor amounts to the pale-yellow Ch forms. Notably, the AHþ and Q species are reversibly interconverted forms by pH changes [175] while, trans- and cis-Ch can in principle be interconverted by photoexcitation. As a result, being the Q and cis-Ch in tautomeric equilibrium, both pH and light stimulation can be used to cause interconversion among the aforesaid fundamental forms (see Fig. 4.12).

Figure 4.12 Anthocyanins equilibria and corresponding structural transformations depending on pH values. In the flavylium cation (red depicted), R3’ and R5’ are H, OH, or OCH3; R5 is OH or O-glicoside moiety (O-Gly).

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It is worth to emphasize that the deprotonation of AHþ cation to form the Q-base is significantly faster (ms timescale) than that of hydration rate toward the H species (s timescale), the latter being faster than that of the opening ring leading to Ch (from minutes to hours). However, because the hydration equilibrium constant, kh, is greater than that of proton transfer equilibrium, kp, the Q (and/or Q) base are unstable and, depending on the pH of the medium, rearrange (through AHþ) into the H-pseudo base and Ch. This means that Q can be formed as a kinetic product which later, totally or partially, disappears to yield a thermodynamic equilibrium [175e177]. In this respect, at equilibrium the system can be described in terms of a single acidebase equilibrium between the acid AHþ and a conjugate base, ConB, having concentration equal to the sum of the concentration of the other species. In natural vegetables, where pH is expected to depend by their fresh and processed nature, each anthocyanin will be, most probably, a mixture of all the aforesaid species. Although the complexity of the Ants behavior can be rationalized within the framework shown in Fig. 4.12, it also is well recognized that the pH-dependent mole fraction distribution at equilibrium (thermodynamic) as well as the rate conversion (kinetics) among the different species, will be different depending on the substitution pattern of the flavylium backbone. As an example, in 3-substituted flavylium compounds as well as in most common anthocyanins, the prevalent product of the ConB is the H-pseudo base, and the Ch-isomers are minority. Differently, 4-substituted flavylium compounds generally equilibrate between the AHþ and Q base, which is also the major species at the equilibrium in 5substituted flavylium or deoxyanthocyanins [175]. Depending on pH, it is also possible the trans-Ch formation even though eOH groups in position 3 make anthocyanins unstable, decomposing into irreversible products [178]. 4.4.3.2 Photochemical and photophysical behavior The possibility of tailoring the electronic and optical properties of Ants is one of the most intriguing and appealing aspects which justifies the widespread interest for this class of compounds. The extended p-conjugation spread over the entire molecule, as well as the presence of a positive charge on the flavylium cation, drives the capability of anthocyanins to absorb strongly in the wavelength region of the UV and visible spectrum with maximum absorption in the 450e550 nm range (band I) and other minor absorption bands around 280 nm (UV

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region, band II). Highly delocalized p-conjugated system in Ants and in natural polyphenols has also been confirmed by quantum chemistry calculations [179,180]. Additionally, the alcohol protons of the Cth ring in the main structure, quickly condense with the eOH groups at the TiO2 nanostructured surface leading to a strong anchorage as well as fast and energetically optimized charge transfer (CT) from Ant to semiconductor CB [181e183]. The anchorage ability of anthocyanidin (Ants without sugar fragment) onto TiO2 surface, through Cth chelating effect, was first reported by Tennakone et al. [184,185], who observed a shift toward higher wavelength (lmax) in the absorption spectrum of “TiO2/cyanidin” system as due to the formation of a stable Ti(IV)-cyanidin complex. Like in synthetic flavylium salt, Ants color and reactivity are affected not only by pH but also by the number and position of the hydroxy (-OH) and/ or methoxy (-OCH3) substitution patterns of the benzo-pyrilium ring. Usually, an increase in the hydroxylation and methoxylation on the B-ring results in a bathochromic shift (shift of the maximum absorption band toward longer wavelengths, l). A comparison among pelargonidin (four eOH groups), cyanidin (five eOH groups), and delphidin (six eOH groups) reveals a lmax, which shifts from 520 to 535 nm and finally 546 nm, respectively (see Table 4.2). Significant effects are also played by the C3 position in the C-ring. For instance, the luteolinidin that have the same number of eOH groups as pelargonidin, but no hydroxyl in C3 position, shows a blueshift of about 25 nm with respect to pelargonidin [175]. It is also suggested that the eOH group in the para-position of the B-ring (C4’ position) increases the p-electrons delocalization in the chromophore [186]. The increasing number of eOH groups in B-ring, that is known to mostly contain the LUMO electron density of Ants [187], decreases the HOMOe LUMO gap allowing photons of longer ls (i.e., smaller energy) to be absorbed, a feature due to a stronger electron affinity of the substituted oxygen compared to hydrogen relative to the sp2 hybridized carbons [188]. In addition, the auxochromes eOH groups, absorbing photons with their nonbonding electrons (n / n*), promote the absorption of lower energies [189]. Furthermore, the shift toward longest ls due to the eOH as well as eOCH3 groups, reflects a higher charge-transfer interaction and concomitantly an increase of the molecular electronic moment dipole, m, which plays an important role in the optical properties. Kouari et al. [190] observed that the attachment of eOH as well as eOCH3 groups (electron-donating groups) in delphinidin, petunidin, and malvidin decrease the HOMOe

Anthocyanidin

R3

R3’

R5’

m(D) [190]

DE [190] HOMO-LUMO

Pelargonidin, Pg Cyanidin, Cy Delphinidin, Dp Malvidin, Mn Peonidin, Pn Petunidin, Pt

OH OH OH OH OH

H OH OH OCH3 OCH3 OCH3

H H OH OCH3 H OH

2.65 4.07 5.63 6.84 4.22 6.29

6.26 6.06 5.96 5.94 6.05 5.99

R3 lmax (nm) OH w520 w535 w546 w542 w532 w543

[191] O-gluc w516 w530 w543 w546 w536 w546

Luteolinidin: R3’ ¼ OH, R5’ ¼ R3 ¼ H

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Table 4.2 Structure of the main widespread anthocyanin-3-O glucoside and respective lmax in the visible spectral region.

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LUMO energy splitting gap and correspondingly increase the ground state m, with respect to pelargonidin, cyanidin, and peonidin. In Ants, the flat polarizable chromophores of the colored forms (mainly AHþ) are quite prone to form molecular associations, via p-stacking interaction, with themselves (self-association) and with other electron-rich molecules (copigmentation) such as flavones, flavanols, organic acids, and so on [192]. Generally, such kinds of association enhance the anthocyanin optical density (hyperchromic effect) and may also result in a bathochromic shift in the absorption spectra (UV-vis region) as well as in a stabilization of the AHþ form boosting the proportion of anthocyanin molecules in the red-colored form [193,194]. Depending on interacting species, intra- or inter-molecular pathways may take place. When the copigment is another phenolic compound, the lack of a real chemical bond makes the interaction transitory. The interaction takes place between compounds with opposite charge, and the electronic density is transferred from the richer to the poorer ring in such a way that the resulting rings are linked by a weak bond. In this regard, the positive AHþ ion results a suitable candidate to form CT-complexes, via pep, with rich electrons compounds. The consequence of the resulting pep complex will be the displacement of the hydration equilibrium (see Fig. 4.12) toward the AHþ form with an increase of its absorption intensity and lmax. Contrarily, when the interacting species have free electron pairs, an intermolecular copigmentation follows and the copigment, covalently linked to the Ants, becoming part of the molecule [195]. Metal cation (e.g., Mg2þ, Sn2þ, Cu2þ, Fe2þ, Al3þ, etc. [192]) can also participate in the copigmentation reaction with anthocyanidin and Ants having o-di-hydroxyl groups in their B-ring (Cy, Dp, Pt) to form metaleanthocyanin complexes. The blue color stability of many flower and plants has been ascribed to the existence of Antemetal complexes [196e198]. Glycosylation and acylation are other critical factors that influence the Ants stability as well as their absorption spectra [199,200]. Glycosylation increase the anthocyanidins stability, and diglucosides are more stable than their corresponding monoglycoside [201]. Introduction of sugar moiety (usually at eOH groups in C3 and/or C5 position) shifts the visible absorption band toward shorter wavelengths (hypsochromic or blueshift) with an appreciable absorption decrease reflected by a lowering color intensity. The absorption spectra of malvidin-3-O-glucoside with that of the corresponding aglycone are compared in Fig. 4.13A, as typical. Possible explanation of this behavior is the enhancement of the flavylium cation electrophilicity due to electrons’ withdrawing effects induced by the sugar moiety [202].

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Figure 4.13 A) Absorption spectra of malvidin-3-O-glucoside antocyanidin and corresponding aglycone form. In the inset, the absorption spectra of delphinidin (three eOH groups) and cyanidin (two-OH groups) aglycones. (B) (b and c) CV and DCV curves of cyanidin and (D) CV of kuromanin (0.1 mmol L1 in MeOH, scan rate 100 mV$s1). (B) and (C) Adapted from M. R. Wasielewski, L. D. Kispert, Direct measurement of the lowest excited singlet state lifetime of all-trans-b-carotene and related carotenoids, Chem. Phys. Lett., 128 (1986) 238e243; T. Gillbro, P. O. Andersson, R. S. H. Liu, A. E. Asato, S. Takaishi, R. J. Cogdell,, Location of the carotenoid 2Ag-state and its role in photosynthesis, Photochem. Photobiol., 57 (1993) 44e48.

It was also suggested that the sugar moieties could cause a decrease in conjugation because of a loss of coplanarity of ring B with respect to the rest of the molecule [203]. As an example, cyanidin and pelargonidin (where there are not sugar moiety) preserve the molecular structural coplanarity maintaining the B ring in the same plane of benzopyrylium through hydrogen interaction between the eOH group and hydrogen atom at positions C3 and C6’, respectively [204]. The influence of substituents (i.e., glucoside, eOH, or eOCH3 groups) on the electrochemical properties of Ants and their corresponding aglycones can be detected by cyclic voltammetry (CV) measurements [205], where the oxidation potential peaks in the corresponding are directly associated to the substituents’ electron-donating power. Fig. 4.13 BeD compares, as demonstrative, the voltammogram curves of two molecules (cyanidin and kuromanin) having the same Cht in the B-ring but different substituents at the molecule C3 position. Cyanidin curve exhibits four peaks (more clearly discernible in the differential pulse voltammetry, DCV curve (Fig. 4.13C) corresponding to a well-defined oxidizable center in the molecule: the two fewer positive assigned to the oxidation of eOH groups at positions C3’ and C4’, the other progressively more positive consistent with the C3,

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C5, and C7 positions, respectively. By a comparison with kuromanin, having the glucoside moiety at C3 position, differently from cyanidin bearing an eOH group, a shift of all oxidation peaks toward more positive, is clearly discernible. Critical effect in determining the anthocyanins spectral behavior (hypsochrome or bathochromic shift) is also played by solvent polarity (solvatochromism). Usually, hypsochrome shift (blueshift) of a chromophore dissolved in solvents with increasing polarity correspond to a negative solvatochromism, while dissolution in solvents with decreasing polarity produces bathochromic shifts and positive (redshift) [206]. In 2014, Meghea et al. [207] reported for pelargonidin-3-O-glucoside dissolved in a binary mixture of water/alcohol, a bathochromic shift inversely proportional to the value of an empirical parameter correlated to the solvent polarity: the lmax absorption band increases with decreasing solvent polarity in a rather linear manner. Finally, cosensitization has been explored as a promising route to extend the spectral response over a broader wavelength region [208e210], enhancing the device current output and thus, the overall efficiency. By cosensitization technique, the usage of two different sensitizers, a more efficient light absorption is expected when a first shorter wavelength absorbing dye is placed in series with a second longer wavelength absorbing dye. The narrowness of the spectral response of a single dye limits the photocurrent of such cells, since it does not cover the full solar spectrum.

4.4.4 Carotenoids Carotenoid polyenes exhibit intense absorption bands in the 300 e 500 nm wavelength range. They play a prominent role supporting Chls in singlet and triplet energy transfer processes and mediate (antenna effect) the dissipation of excitation energy absorbed in excess from the photosynthetic system [211e213]. Generally, when in solution carotenoids are photochemically not reactive; however, they can undergo unimolecular isomerization and act as quenchers of the triplet states of other compounds through energy transfer. This probably depends on the very short lifetime of their singlet excited state (50 ps < s < 10 ps) [214,215]. In 2000, Gao and coworkers [216] published their pioneering study on carotenoids (80 -apo-p-caroten80 -oic acid (ACOA)) as sensitizer for TiO2 nanocrystalline mesoporous electrode. Despite the very low Voc and Jsc values achieved by their photoelectrochemical device, their work was a breakthrough to overcome the

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kinetic limitations for efficient electron photoinjection and carrier transport imposed by the short life of single excited state of carotenoids, through the introduction of terminal carboxylate groups and the generation of covalent bonds with the TiO2 surface. In 2010, Gomez-Ortiz and coworkers [217] explored purified carotenoids, bixin and norbixin, from the extract of achiote seeds (annatto [218]), as sensitizers for DSSC applications with the best results reached in bixin-sensitized TiO2 solar cell. Achiote is an American continent shrub that bears an inedible fruit containing red seeds from which a dark-red extract (annatto) can be obtained. The annatto seeds pericarp contains mainly cis-bixin (methyl hydrogen (90 Z)-6,60 -apocarotene-6,60 -dioate, molecular formula C25H30O4, up to 80%) and trans- and cis-norbixin (6,6’-diapocarotene-6,60 -dioic acid and 90 cis-6,60 -diapocarotene-6,60 -dioic acid, molecular formula C24H28O4, up to 20%) (Fig. 4.14). Cis-bixin is a water-insoluble pigment, constitutes by a long chain of alternating double conjugated bonds, with carboxylic acid and methyl-ester groups each at each end of the chain, while trans- and cis-norbixin with a carboxylic acid moiety instead of methyl ester moiety is water soluble. In bixin, the carboxylic group at C7 carbon atom in a conjugated long chain, being not methyl carboxy ester, easily binds with the

Figure 4.14 A) Absorption spectra of annatto-1 solution extracted by diethyl ether and annatto-2 extracted by acetone. The inset shows the bixin spectrum. (B) Absorption spectra of TiO2 films sensitized by annatto-1 and annatto-2. In the inset of panel (B) the photoaction spectrum of dye-sensitized solar cell sensitized by annatto-1 together with the bixin and norbixin chemical structures are also reported. Adapted from G. Calogero, J. Barichello, I. Citro, P. Mariani, L. Vesce, A. Bartolotta, A. Di Carlo, G. Di Marco, Photoelectrochemical and spectrophotometric studies on dye-sensitized solar cells (DSCs) and stable modules (DSCMs) based on natural apocarotenoids pigments, Dyes Pigments 155 (2018) 75e83.

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hydroxyl groups at the TiO2 surface in an effective monodentate linkage. In addition, both bixin and norbixin show ε as high as 1.9∙105 Lmol1∙cm1 and 1.4∙104 Lmol1∙cm1, respectively, and consequently high lightharvesting ability, bixin being the best. In this framework, our research group in collaboration with CHOSE researchers explored conditions under which annatto pigments from achiote seeds can elicit new interest as proper sensitizers in TiO2-nanostructured photoanode to generate photocurrents in DSSCs and in dye-sensitized solar modules (DSCMs). It was demonstrated via UV-Vis spectrophotometric investigations that the use of less polar solvent such as diethyl ether, for the dye’s extraction process, is preferable than that of ethyl acetate or acetone [219]. From this novel extracting-solvent approach, a dark-red extract (annatto-1) was employed to sensitize a constructed cell/module. In addition, by examining the influence of TBP electrolyte additive in the apocarotenoid-based DSCs, TBP-free electrolytes were revealed as the better choice for this kind of sensitizers. The constructed device showed an IPCE maximum of 32.8% with an estimated Finj  0.44. Also, by optimizing the TiO2 layer thickness in the 6e10 mm range, a gain of w182% and w33% in Jsc and VOC, respectively, was achieved. Finally, by comparing the obtained results (h ¼ 1.6%, Jsc ¼ 8.52 mA/cm2 for annatto-1 sensitizer) with previous literature data, there was an impressive increment of h (w742%) and Jsc (w1507%) values. When apocarotenoids extracted from the achiote’s seeds were used in DSCM, a power of w8 mW, with VOC of w1.7 V and current Isc of w13 mA, was measured. Furthermore, shelf life stability (1000 h) under 1 sun (AM1.5) was reached with an estimated battery capacity of w46.8 Ah (corresponding to 15 standard type AAA alkaline batteries). The experimental results, as reported in Figs. 4.14 and 4.15, open new expectations toward long-term stable vegetable pigmentebased solar modules for electric purposes.

4.5 Conclusion The main aim of this chapter was a comprehensive discussion on natural sensitizers to provide low-cost, environmentally sustainable, and easy to fabricate novel DSSCs. Explanations to highlight the DSSC figures of merit and working principles with respect to the more conventional PV devices have been given. Emphasis has also been devoted to the photosensitizer structure and their anchoring groups, essential to develop cheap alternative to Ru/organic metal-free complex dyes in terms of intrinsic stability, charge

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Figure 4.15 Currentevoltage characteristic of ye-sensitized solar module (DSCM) sensitized by annatto-1. A picture of the constructed DSCM compared to a 2V coin together with the short-circuit current (ISC) and efficiency h along 1000 h of shelf life are also reported in the insets. The electrical performance of the device was measured under a Class B Sun Simulator (Solar Constant 1200 KHS) at AM 1.5G, 100 mW cm2, using a Keithley 2420 as a source meter in ambient conditions. Adapted from G. Calogero, J. Barichello, I. Citro, P. Mariani, L. Vesce, A. Bartolotta, A. Di Carlo, G. Di Marco, Photoelectrochemical and spectrophotometric studies on dye-sensitized solar cells (DSCs) and stable modules (DSCMs) based on natural apocarotenoids pigments, Dyes Pigments 155 (2018) 75e83.

transport capability, and high compatibility among the cell standard components. By comparing synthetic with natural dyeebased DSSCs, we need to consider the enormous efforts made to develop efficient synthetic dyes [45], which are still lacking for the natural counterpart. In this context, many issues have to be overcome before this technology can become truly commercially for large-scale implementations. Efficient DSSCs require intense light absorption and rapid charge injection from the dye to the semiconductor NPs. Usually, the low VOC value, affecting the majority of natural dyeebased DSSCs, is due to inefficient reduction of the oxidized dye and/or adsorption of environmental Hþ ions, causing a positive shift of the TiO2 Fermi level and thus limiting the maximum VOC achievable by the cell. Also, the steric hindrances and dye pigment molecules aggregation limit its ability to bind uniformly the semiconductor surface as well as its efficient absorption onto the TiO2 NPs film,

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respectively, thus affecting both light harvesting and charge transfer, which in turn determine the cell low efficiency. Several kinds of additives and coabsorbers have been explored to reduce the aggregation effect. However, in several pH-sensitive natural dyes (e.g., Ants or betalains), the use of basic additives (e.g., TBP and similar), tough able to increase the VOC value, on the other side, influences the dye ε, and the resulting shift in the absorption maximum toward the spectral red region reduces the dye-sensitized activity. The natural dyes LUMO orbital is pushed to more negative values, while the TiO2 CB moves toward less negative potential, resulting in a weak driving force and a small electron injection. Natural dyes can strongly contribute to reduction of environmental impact, and also represent an interesting and multidisciplinary research activity useful for dissemination of knowledge and to educate people on renewable energy sources. Research studies are going on for improving the efficiency and viability of natural dye-based DSSCs, including solid-state electrolytes for better temperature response, lifetimes, and degradation dye mechanisms. The cell design, too, may be improved in order to put this novel technology in position to compete with the more conventional counterpart. The opportunities to utilize sunlight as renewable and inexhaustible energy source evolve dynamically, as a function of human ingenuity.

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CHAPTER FIVE

Perovskite solar cells Narges Yaghoobi Nia1, * Danila Saranin2, Alessandro Lorenzo Palma1, 3, Aldo Di Carlo1, 2 1

Centre for Hybrid and Organic Solar Energy, University of Rome Tor Vergata, Rome, Italy Laboratory for Advanced Solar Energy, National University of Science and Technology ‘‘MISiS’’, Moscow, Russia 3 Italian National Agency for New Technologies, Energy and Sustainable Economic Development (ENEA), Energy Efficiency Unit Department, Rome, Italy *Corresponding author. 2

5.1 Introduction Hybrid organiceinorganic halide perovskite solar cells (PSCs) have risen to stardom owing to the peculiar characteristics of the halide perovskite absorber such as high charge carrier mobility, broad and strong optical absorption, long free carrier diffusion length, low exciton binding energy, as well as their cost-effective and easy solution process manufacture [1,2]. PSCs have achieved remarkable progress with power conversion efficiency (PCE), developing from 3.8% in 2009 to 23.7% in 2018 [3,4]. PSCs have become an important technology in the photovoltaic (PV) field, rivaling the well-consolidated silicon-based solar cells, copper indium gallium selenide (CIGS) and cadmium telluride (CdTe) solar cells [5]. This excellent improvement in PCEs is largely due to the optimization of the device’s architecture, the use of interface engineering, improvement of hole and electron transport materials, and the development of fabrication processes with high-quality perovskite films [9e14]. These progresses have promptly motivated the PV communities’ efforts to commercialize PSCs. Many efforts have been focused to scale up PSCs, and great progress has been achieved in perovskite module performance [5e7]. Many excellent review papers have been published to summarize the advances of upscaling of PSC [15e19]. However, considering that large-scale manufacturing of perovskite films is mainly based on solution processes, an in-depth understanding of nucleation and growth processes of perovskite crystals is essential to obtain a high-quality perovskite film, which is directly linked with the performance of PSCs.

Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00005-7

© 2020 Elsevier Ltd. All rights reserved.

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In this chapter, we first summarize the main properties of halide perovskite, and then we review the mechanism of nucleation and growth of perovskite thin films. PSC architectures and scaling up procedures are then discussed in the last two sections of the chapter.

5.2 Unique properties of metaldhalide perovskites for photovoltaics 5.2.1 Molecular composition and basic materials PSCs are based on an absorber with ABX3 perovskite crystal structure, where “A” and “B” are cation elements and “X” is an anion element. Originally, the terminology of “perovskite” materials was accepted in the first half of the 19th century, when the group of Lev Perovskii and Gustav Rose discovered mineral-CaTiO3 with ABX3 structure. Currently, the family of perovskite materials has a lot of fields for applications like superconductors [20,21], optical waveguides [22], transparent electrodes [23,24], etc. Mainly based on oxides, they achieved progress and lasting successes in PV in the form of hybrid organiceinorganic halide perovskites. The organometallic composition for ABX3 crystal (schematically presented in Fig. 5.1) has showed the most advantageous semiconductor properties for use in thin film solar cells [25]. “A”dsite represents small organic cations such as methyl ammonium (MAþ, CH3 NH3 þ ); formamidinium (FAþ, CH4 N2 þ ); guanidinium (GUAþ, C3 N3 H6 þ ), as well as a single element from the first group of periodic table (PT) such as Csþ, Kþ, Rbþ. “B”dsite cation represents an element from the 14th group of the PT such as Pb, Sn, and Ge. “X”dsite anions represent halogens of the 17th group such as I; Br, and Cl. Appropriate formulation of A, B, and X components should satisfy the Goldschmidt tolerance factor [26], a specific ratio between the ionic radius of cations and anions, described in Eq. (5.1): rAeff þ rX t ¼ pffiffiffi 2ðrBþ rX Þ

(5.1)

where rA,B,X is the effective ion radius of A, B, and X. t should be in a range between 0.80 and 1.15 [27] to form an organometallic semiconducting perovskite crystal. The range of values between 0.8 and 1 characterizes perfect fitting between A-cation and BX6 octahedrons, while the structures with smaller t (around 0.8) has increased tilting of the BX6 that tends to form lower symmetry crystals. If the ionic radius of A is too big (t > 1), a more

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Figure 5.1 The ideal cubic perovskite unit cell. (A) A cations(blue) occupy the lattice corners, B cations (green) occupy the interstitial site, and X anions (red) occupy lattice faces. (B) An alternative view depicting B cations assembled around X anions to form BX6 octahedra, as B X bonds are responsible for determining electrical properties. (C) Tilting of BX6 octahedra occurring from nonideal size effects and other factors, inducing strain on the B X bonds. Reprinted from Q. Chen, N. De Marco, Y. Yang, T. Bin Song, C.C. Chen, H. Zhao, Z. Hong, H. Zhou, Y. Yang, Under the spotlight: the organicinorganic hybrid halide perovskite for optoelectronic applications, Nano Today. 10 (2015) 355e396. https://doi.org/10.1016/j.nantod.2015.04.009

complex structure is formed with intermolecular distortions. Typically, if t < 0.8, A-B-X ions do not form perovskite crystals. In opposite to inorganic oxide compositions, metaleorganic halide perovskites do not have spherical symmetry of the organic cation. In this case, the correct estimation of organic cation size becomes more complicated, as well as a precise calculation of tolerance factor with accounting of cationeanion interactions. Thus, Goldschmidt tolerance factor is not a comprehensive parameter that completely describes stability [28] of the crystal structures (Fig. 5.1) that can be formed between A cation and the corner-sharing BX6 octahedra in dependence to their size. For this reason, a special geometric octahedral factor, m [29], was also assessed for determination of tilting between B and X sites. An objective approach for the prediction of ABX3 stability requires combinational account of t-m parameters. The ionic size of cation and anions in A-B-X sites influences the ordering/disordering of the molecular symmetry and can shift the orientation alignment, off-centering, and bonding [30]. Thus, phase transformation of metalorganic halide perovskites strongly depends on the stoichiometry and composition. Moreover, organic parts (representing A site cation) add functionality for interaction with inorganic BX6 parts. Hybrid halide perovskite crystals have several phase transitions in 100e400 K temperature range [31]. For example, for MAPbI3 crystals [32], reorientation of A site

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Figure 5.2 Octahedra tilting pattern for the tetragonal (left) to the cubic (right) structure of a MAPbX3 perovskite (X ¼ halide); the order parameter of the transition. Reprinted from C. Quarti, E. Mosconi, J.M. Ball, V. D’Innocenzo, C. Tao, S. Pathak, H.J. Snaith, A. Petrozza, F. De Angelis, Structural and optical properties of methylammonium lead iodide across the tetragonal to cubic phase transition: implications for perovskite solar cells, Energy Environ. Sci. (2016). https://doi.org/10.1039/c5ee02925b.

CH3NHþ 3 first cause transition from Pnma to tetragonal I4/mcm symmetry at w170 K, then to photoactive, cubic Pm3m phase at w330 K (see Fig. 5.2), and finally to P4mm at a temperature >400 K with large cation disordering. The use of mono FAþ or Csþ cations can lift up the transition temperature to “black” photoactive phase at temperature >420 K [33] and >570 K [34], respectively. The most studied perovskite composition for prototypes of PSCs is MAPbI3, due to a low temperature crystallization (100 C) and high solar cell performance (>20%). However, due to a weak alignment between MA and PbI6 sites (tfactor ¼ 0.83), low temperature transitions into cubic photoactive phase (w320 K) occur and cause several instabilities [35], which will be discussed in the next subchapters. Actual trends for perovskite molecule stabilization include double/triple cation compositions to balance the molecular packing density.

5.2.2 Band gap structure In general, organometallic perovskites have direct band gap structure (Fig. 5.3A), which minimizes the thermalization losses which occur in nondirect band gap Si [36]. Many simulations and theoretical calculations, mainly based on density functional theory, have been performed to estimate the electronic band structure [36,39,40], and comparisons with experimental data are quite good. Optoelectronic properties of halide perovskites can be easily tuned by replacing cations and anions in the composition. A change in

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Figure 5.3 (A) Calculated (SOC-GW) band structure (VB and CB highlighted as thick solid lines) for MAPbI3 (blue lines). (B) Absorption coefficient of (MA)Pb(BrxI1-x)3 measured by diffuse spectral reflection and transmission measurements on thin films and photocurrent spectroscopy of solar cells. SOC, spineorbit coupling; VB, valence band; CB, conduction band. (A) Reprinted from P. Umari, E. Mosconi, F. De Angelis, Relativistic GW calculations on CH3NH3PbI3 and CH3NH3SnI3 perovskites for solar cell applications, Sci. Rep. 4 (2015) 4467. https://doi.org/10.1038/srep04467. (B) Reprinted from E.T. Hoke, D.J. Slotcavage, E.R. Dohner, A.R. Bowring, H.I. Karunadasa, M.D. McGehee, Reversible photo-induced trap formation in mixed-halide hybrid perovskites for photovoltaics, Chem. Sci. (2015). https://doi.org/10.1039/c4sc03141e.

the anionic composition from Cl to Br (X-sites) tunes the Eg from 3.16 to 2.30 eV [40], and down to 1.55 eV with I anion [41], as shown for MAPbX3 perovskite in Fig. 5.3B. The optimal band gap of perovskite for terrestrial single-junction solar cell approximately lies in the range between 1.2 and 1.62 eV, as it was shown for best-performing devices [42]. The classical calculations of ShockleyeQueisser limit [46,47] show that the maximum solar cell efficiency of 33.4% can be achieved with a 1.34-eV band gap semiconductor, and organometallic halide perovskite allows to fit this value with appropriate chemical composition. Eg can be also tuned by varying the A and B site cations. To narrow the band gap and expand the absorption spectrum into the infrared region, the cation B (typically Pb) can be gradually replaced by tin (up to 1.2 eV) [45]. The tuning of A site cation also can shift the Eg width [46] in comparison to standard MAPbI3 crystal. Replacing MA with FA cation decreases Eg to 1.4 eV [47], while CsPbI3 has an Egw1.70 eV [48]. So, all compositional varieties of the cations and anions can effectively change the optoelectronic properties for different applications of the solar cells as single-junction devices, tandems, and detectors for shortwave optical spectrum of light.

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Figure 5.4 Effective absorption coefficient of a CH3NH3PbI3 perovskite thin film compared with other typical photovoltaic materials, including amorphous silicon (aSi), GaAs, CIGS, CdTe, and crystalline silicon (c-Si), all measured at room temperature. €per, B. Niesen, M. Ledinsky, F.J. Reprinted from S. De Wolf, J. Holovsky, S.J. Moon, P. Lo Haug, J.H. Yum, C. Ballif, Organometallic halide perovskites: sharp optical absorption edge and its relation to photovoltaic performance, J. Phys. Chem. Lett. (2014). https://doi.org/10. 1021/jz500279b.

Halide metalorganic perovskites (MAPbI3, for example) are characterized by a sharp onset and abrupt edge in the absorption coefficient above the band gap (Fig. 5.4) [49], typical of direct band gap materials. Lead halide perovskite crystal, due to the heavy atoms present in the structure, has an interesting characteristic caused by strong spineorbit coupling (SOC) of lead and iodine that provides the splitting of electron states. In detail, splitting occurs for electrons with spins directed up and down along the axis of the wave vectors, and this process is called Rashba effect [50]. Such state in the electronic structure appears in the organometallic perovskites that have no inversion and center of symmetry. SOC creates an effective magnetic field that influences the electron states, removes the degeneration of spins, and splits the edges of valence and conduction bands, that finally result in a slight split of the band gap from direct to nondirect with DEw50 meV [51] (as shown in Fig. 5.5). As it was showed in several works [54e58], Rashba splitting in halide perovskite significantly reduces the rates of radiative recombination (for more than 35%), ensures a high density of states, enhances the optical absorption, and increases the charge carrier lifetimes.

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αC =

εC≠ 2ΔkC

ε

ε+C

ΔkC

ε+

εC– k

ΔkV

V

εC≠

V

ε–

ε≠ V

αV =

εV≠ 2ΔkV

Figure 5.5 Schematic representation of Rashba splitting for parabolic bands, along with relevant interaction parameters, see text for definitions. The valence and conduction bands (VB and CB) in the presence (absence) of Rashba splitting are the solid red and blue (dashed) lines, respectively. Reprinted from E. Mosconi, T. Etienne, F. De Angelis, Rashba band splitting in organohalide lead perovskites: bulk and surface effects, J. Phys. Chem. Lett. (2017). https://doi.org/10.1021/acs.jpclett.7b00328.

5.2.3 Crystal instability While PSCs already have an advantage in efficiency competition with CIGS and CdTe solar cells, stability remains a main problem that delays PSC commercialization. In the list of degradation factors of PSCs, crystal instability is a major obstacle for the long-term device operation. Intrinsic instability is mainly caused by photo and thermal factors that induce chemical and structural degradation. Both degradation processes cause decomposition of ABX3 crystal to organic cations, gaseous forms of anions, BX2 salts, and restored pure element B. Several investigations for the MAPbI3 films showed significant phase and composition changes at temperatures higher than 50 C. As it is shown in the work of Juarez-Perez et al. [57], photo stress of sun illumination and thermal stress in 40e80 C induce three degradation mechanisms: (1) organic cation decomposition: CH3NH2 þ HI (reversible path); (2) formation of gaseous products: NH3 þ CH3I (the irreversible or detrimental path); (3) a reversible cationeanion decomposition: Pb(0) þ I2(g). Reversibility of processes and self-healing of MAPbI3 films after thermal and photo stressing can be obtained in dark conditions and low temperatures [58], but nevertheless decomposition and recrystallization processes left light-activated metastable trap states which accumulate from cycle to cycle, and finally significantly decrease PV performance with nonradiative

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recombination effects. Sequential transition from tetragonal MAPbI3 to trigonal PbI2 under illumination was described in the work of Yang et al. [59], where authors indicated the increase of signal from PbI2 component in the electron diffraction patterns during analysis of thermos-kinetics processes with time. Several approaches have been proposed for the stabilization of MAPbI3 device operation [60], like solar cell encapsulation (to prevent leakage of volatile decomposition products) and selection of chemically stable transport layers, that must be chemically inert with acidic molecules (HI), methylation reagents (CH3I), oxidizing agents (I2), and weak bases (CH3NH2, NH3). Nevertheless, MAPbI3 absorbers do not show longterm stability, due to cycling transition processes which spoil the morphology and increase the level of bulk recombination on defectdtrap centers. For this reason, cation engineering should be applied to increase the tolerance factor by using larger cation size and to increase the temperature of transition to cubic phase, to prevent unfavorable thermal decomposition processes at the typical temperatures of the solar cell operation (w350 K). FAþ cation with ionic radius of 2.79 Å (vs. 2.70 Å for MAþ) provides higher temperature of black phase formation w420 K, but, at the same time, single-cation FAPbI3 is not completely stable at room temperature, due to the formation of both hexagonal d-phase (“yellow phase”) and photoactive perovskite a-phase [61]. Inorganic CsPbI3 also forms yellow orthorhombic phase at room temperature and becomes stable only near w300 C [62]. So, it is possible to conclude that monocation halide perovskites cannot perform stable operation because of low-temperature phase transitions. Thus, precision mixing of the cations and halides should result in perovskite thermal and structural stabilization. A big step forward for the reduction of perovskite instability was done with cesium-containing triple cation engineering [63]. Authors showed that double MA/FA compositions still have intrinsic thermal instability that tends to form the less crystalline films. Incorporation of inorganic cesium with small ionic radius (1.81 Å) to A-site suppresses transition processes to yellow phase and improves morphology. A composition of Cs5(MA0.17FA0.83)0.95Pb(I0.83Br0.17)3 was found to be optimal for the performance and stability, with achievement of PCE up to 21.1% and outputs of 18%, even after 250 h of power point tracking (Fig. 5.6). Another approach for the crystal stabilization was realized by Grancini and coworkers [64] with 2De3D modification of perovskite submicron grain structure using large AVAI (HOOC(CH2)4NH3I) cation (Fig. 5.7) that results in bottom-up phase-segregated structures with combination of

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Figure 5.6 Improvement in the long-term device stability with Cs cation incorporation. PCE, power conversion efficiency. Reprinted from M. Saliba, T. Matsui, J.-Y. Seo, K. Domanski, J.-P. Correa-Baena, M.K. Nazeeruddin, S.M. Zakeeruddin, W. Tress, A. Abate, A. €tzel, Cesium-containing triple cation perovskite solar cells: improved staHagfeldt, M. Gra bility, reproducibility and high efficiency, Energy Environ. Sci. 9 (2016) 1989e1997. https:// doi.org/10.1039/C5EE03874J.

Figure 5.7 (A) Local density of state of the 3D/2D interface. (B) Interface structure with the 2D phase contacting the electron transport layer. Reprinted from I. Zimmermann, E. Mosconi, X. Lee, D. Martineau, S. Narbey, F. Oswald, G. Grancini, C. Rolda, One-Year stable perovskite solar cells by 2D/3D interface engineering, Nat. Commun. (2017) 1e8. https:// doi.org/10.1038/ncomms15684.

the 2D layer acting as a protective buffer against moisture, preserving the 3D perovskite of phase transition caused by thermal stress. The use of advanced mesoscopic structure with ZrO2 scaffold and stable carbon cathodes resulted in the most stable configuration of the perovskite device to date, with >10,000 h of stable efficiency under light soaking.

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5.2.4 Charge transport Charge transport is one of the most critical topics for metal halide organometallic perovskites. The core of solar cells functionality lies in the electronic structure and charge carrier conditions. As it was noted before, halide perovskites are direct semiconductors; moreover, tetragonal and “black” cubic phases both have similar positions for valence band maximum and conductive band minimum in the Brillouin zone, so the optical transitions of lightexcited charge carriers will occur without phonon interactions [65]. For MAPbI3 perovskite, the generated electronehole (eeh) couples has binding energy (Eb) around 0.055 eV [66], thus charge splitting does not require high-electric fields across the absorber. As it was found in the literature, basic charge transport characteristics of MAPbI3 microcrystals [65] have a diffusion length (LD) up to 1 mm, a carrier lifetime up to 1 ms [67], and a charge carrier mobility (m) up to 101 cm2/(V*s) [67]. It is important to point out that, differently to classic semiconductors, the listed charge transport parameters approximately apply to both electrons and holes. The origin of such behavior is mainly related to the electronic structures of conduction and valence band. The electronic structure of the metaleorganic halide perovskite presents similar curvature for the bottom of the conduction band and the top of the valence band, consequently the electron and hole effective masses are equal, m*e ¼ m*hw(0.1e0.15)m0 [68, 71], where m0 is a free electron mass. Such values for MAPbI3 polycrystalline films are comparable with conventional inorganic semiconductors for optoelectronic application such as GaAs, InP, etc.

5.2.5 Bulk recombination There are two main types of recombination occurring in solar cells: radiative and nonradiative. In case of PSCs, it was shown [69] that radiative recombination of free electrons and holes in MAPbI3 cannot give a serious negative impact for solar cells operation because of low bimolecular charge recombination constants, high mobilities (>10 cm2/(V*s), and large diffusion lengths (up to 1 mm). On the other hand, nonradiative, trap-assisted recombination (NRR) is a major mechanism for losses in PSCs [70]. Trap centers in perovskite thin films with submicron crystallinity are mainly concentrated at grain boundaries in the bulk due to favorable energy states and at the interface between perovskite and transport materials (Fig. 5.8A and B). In general, intrinsic defects of the halide perovskite are represented by vacancies of anions, organic and metal cations; interstitials, antisite

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(A)

(B) FTO

Ec G

Re

TiO2

Steady-State Model

Er Rh

Ef

Electron Hole Trap

Case 2 Case 1 Case 3

CH3NH3Pbl3 spiro-OMeTAD

Ev

Au

X

Figure 5.8 (A) Schematic energy level diagram of a single-effective recombination center model, representing the carrier generation by illumination and the processes of electron and hole capture in the recombination centers. (B) Device structure of perovskite solar with illustrated recombination processes: Case 1: bulk recombination at grain boundaries, denoted as “Bulk”; Case 2: electron transport layer/perovskite interfacial recombination at its interface, denoted as “Top”; Case 3: perovskite/hole transport layer interfacial recombination at its interface, denoted as “Bottom.” (A) Reprinted from I. Levine, S. Gupta, T.M. Brenner, D. Azulay, O. Millo, G. Hodes, D. Cahen, I. Balberg, Mobilityelifetime products in MAPbI 3 films, J. Phys. Chem. Lett. 7 (2016) 5219e5226. https://doi.org/10.1021/acs.jpclett.6b02287. (B) Reprinted from L.X. Shi, Z.S. Wang, Z. Huang, W.E.I. Sha, H. Wang, Z. Zhou, The effects of interfacial recombination and injection barrier on the electrical characteristics of perovskite solar cells, AIP Adv. (2018). https://doi. org/10.1063/1.5021293.

occupations, that give a main contribution to NRR, for example, vacancies of the anion sites in MAPbI3, have relatively low energy of activation in the range of 0.08e0.90 eV and can be electrically mobile [75,76]. Measurements of temperature-dependent capacitance frequency for MAPbI3 films demonstrated that shallow traps are present with an energy level of around 0.20 eV from the conduction band, and deep traps of w0.50e1.1 eV can be obtained with concentrations of 1015ee1016cm3 [78e80]. Actual efforts for grain boundaries passivation are based on advanced methods of crystallization as well as incorporation of special additives for morphology modification. Niu and coauthors [76] showed that addition of DMSO to perovskite solution improves the coordination of Pb2þ during nucleation and passivate antisite PbeI states via more heterogeneous growth of perovskite frameworks. As a result, PCE was improved from 16.5% to 18.5%. Ye et al. [77] proposed the incorporation of Cu(thiourea) I to the MAPbI3 absorber for coordination of metal cations and halide anions with a passivation agent at grain surface. Such intergrain passivation improved transient PL spectra with delayed quenching, provided increased Voc for the devices from 1.00 to 1.11 V, and boosted the efficiency from 8.5% to 18.6%.

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Lee and coworkers presented a concept of 2D/3D crystal structure formation with the addition of phenethylammonium cation (PEAþ) [78]. A small amount of additive (15%) without the need for complex mesostructures (Fig. 5.18B). In order to follow the concept of organic PV devices with HTL on transparent electrode, Guo et al. [120] demonstrated the first peien oriented planar solar cells (Fig. 5.18C) with organic transport layers PEDOT:PSS (poly(3,4ethylenedioxythiophene) polystyrene sulfonate) and PCBM (Phenyl-C61butyric acid methyl ester). The unprecedented progress in the efficiency growth of PSCs has been achieved with a scientific supply of experience from the development or DSSC and OPV. Even so, halide perovskite rapidly became a separate class of semiconductor technology for PV, with their own specifics of device physics. The actual classification of PSC architectures can be divided in two main groups accordingly to the orientation of transport layers: “direct” structure with n-i-p orientation and “inverted” structure with p-i-n sequence. These groups can be further divided into mesoscopic and planar configurations (see Fig. 5.18). Efficiencies of described device architectures can exceed 20% PCE by using optimized perovskite crystallization methods, performing transport layers and passivation techniques [130e132]. So far, all the PSC records

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[124] have been achieved by using the direct mesoscopic structure, while planar configurations are slightly below: neiep planar structure recently demonstrated 21.6% [125] and inverted peien structure showed 20.91% [126]. PSC structures can be also realized in pen heterostructure configuration (FTO/TiO2/perovskite/metal) as presented by Etgar [127]; however, the performance of such architectures is not competitive with neiep or pe ien structures. The variety of transport layers including numerous quantities of inorganic and organic materials [137e140] is presented in Fig. 5.19 (see energy levels for typical transport layers). In principle, a charge transport layer should provide perfect level alignment with perovskite’s conduction and valence band for the collection of hþ and e, respectively (see schematic band diagram on Fig. 5.19). The absence of ohmic losses at heterojunctions will provide high values of Voc, since its value is related to the difference between quasi-Fermi levels (QFLs). For ideal conditions, in PSC the QFLs are constant across the perovskite film and the difference between them is the open circuit voltage (Fig. 5.20)

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Figure 5.19 Energy level diagrams with respect to vacuum (0 eV) for typical hole transport layer (HTL) and electron transport layer (ETL) with mono-cation perovskites. Reprinted from V. Zardetto, B.L. Williams, A. Perrotta, F. Di Giacomo, M.A. Verheijen, R. Andriessen, W.M.M. Kessels, M. Creatore, Atomic layer deposition for perovskite solar cells: research status, opportunities and challenges, Sustain. Energy Fuels. (2017). https://doi.org/ 10.1039/c6se00076b.

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Figure 5.20 Band diagram of an ideal solar cell under open circuit conditions, with the absorber.

[132]. As noted before, diffusion lengths in perovskite for both types of charges can exceed the thickness of the perovskite film. In this case, the open circuit voltage can be affected by nonradiative recombination in the bulk and at the interfaces with transport layers.

5.4.2 Tandem solar cells with perovskites The development of tandem solar cells attracts the interest of the scientific community as an effective approach to overcome the ShockeyeQueisser limit, formulated for ideal single cell with 1.34 eV band gap and maximum efficiency of 33.7% [133]. Multijunction solar cells with aligned spectral splitting of the subcells already reached high PCE 32.8% for two-junction and 37.9% for three-junction devices, as reported for AIIIBV semiconductors. Yet, unfortunately, manufacturing of large-scale AIIIBV devices is still an expensive and complex technology for mass production. Progress in the increase of the output performance of silicon and CIGS reached the saturation at 26% and 23% PCE without concentration of light [124]. The rapid growth of PSC efficiency demonstrated that this novel class of semiconductors has a big potential for tandem integration with commercially available Si and CIGS solar cells, due to a comparable level of performance and easy manufacture. Moreover, cost estimates and economic analyses performed for PSCs showed that the expected price for a 1-m2 module will be w30 US$, dominated by the cost of glass substrates and back cover (10 US$). The cost of absorber, transport materials, contacts, and TCOs does not exceed 5 US$ per m2, other part of the prices going for packing, lamination, etc. [144e146]. In this case, integration of PSCs with silicon solar cells that present cost of 70e80 US$/m2 will not significantly increase the cost of fabrication [136].

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Figure 5.21 Schematic illustration of spectral splitting in tandem solar cells with decrease of photons energy absorbed in bottom subcells. Reprinted from S. Kurtz, J. Geisz, Multijunction solar cells for conversion of concentrated sunlight to electricity, Optic Express 18 (2010) A73. https://doi.org/10.1364/OE.18.000A73.

Opposite to single-junction solar cells, tandem configuration allows to minimize thermalization losses with light absorption via spectral splitting in subcells with different Eg. In general, tandem architecture has a stack of subcells with large band gap on the top and narrow band gap on the bottom. The top cell is transparent to low-energy photons that go through the device and are absorbed in the bottom cell (see Fig. 5.21). Tandem devices can be realized in two-terminal (2T) and fourterminal (4T) configuration (Fig. 5.22). The 2T configuration requires a series connection between the cells and consequently a current matching between them. On the opposite, 4T tandem does not require current matching, but introduces an increased level of parasitic absorption and reflection at the electrodes between the subcells and an increased PV system cost due to the request of independent maximum power point tracker. Simulations of the theoretical tandem efficiency considering (1) a 100% EQE above the band gap, (2) a reverse saturation current calculated assuming 100% radiative emission, and (3) 1.1 eV band gap for the bottom subcell (representing Si and CIGS SC) [136] showed that the highest PCE can be obtained with a top perovskite subcell with Eg ¼ 1.60e1.75 eV, resulting in a predicted performance of 44%.

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Figure 5.22 Device schematics of two-terminal (A) and four-terminal tandems (B) on the example of perovskiteeSi structures. Reprinted from J. Werner, B. Niesen, C. Ballif, Perovskite/Silicon Tandem Solar Cells: Marriage of Convenience or True Love Story?-An Overview, (2017). https://doi.org/10.1002/admi.201700731.

Present record for perovskite/Si 2T tandem shows a PCE ¼ 28.0%, achieved by Oxford Photovoltaics [133,149], while perovskite/CIGS tandem developed by IMEC (Solliance) gave PCE ¼ 24.60% [138], but the details about spectral matching and the compositions of the perovskite subcell were not disclosed. The most efficient perovskite/Si tandems were designed with multication perovskite compositions with 1.63e1.74 eV band gaps. Bush et al. [139] developed tandem Si heterojunction solar cell with double cation CsFAPbI3xBrx perovskite absorber with 1.63 eV band gap; the device structure was monolithically integrated and balanced at 18.5 mA/cm2 current density, and finally a 23.6% PCE was achieved with 1000-hour damp heat test at 85  C and 85% relative humidity (Fig. 5.23). Peng et al. [140] used triple cation Cs0.05(MA0.17FA0.83)0.95 Pb(I0.83Br0.17)3 with Eg ¼ 1.63 eV in 4T mechanically stacked tandem with two ITO electrodes in PSC structure with IBC silicon cell and obtained a 24.5% PCE. Duong et al. [141] developed four-cation RbCsMAFAPbI3xBrx composition of perovskite absorber with Eg ¼ 1.74 eV in same device stack and achieved 26.4% PCE via reduced optical losses in transport layers (with use of wide band gap TiOx-doped with In doping), improved transmittance in visible region and decreased contact resistance. In all cited papers, the efficiency of subcell was exceeded by 1%e3% when used in tandem structure with perovskite.

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Figure 5.23 Device schematic of monolithically integrated perovskite-Si tandem (A), balanced spectral splitting between Si and perovskite subcells. Reprinted from K.A. Bush, A.F. Palmstrom, Z.J. Yu, M. Boccard, R. Cheacharoen, J.P. Mailoa, D.P. McMeekin, R.L.Z. Hoye, C.D. Bailie, T. Leijtens, I.M. Peters, M.C. Minichetti, N. Rolston, R. Prasanna, S. Sofia, D. Harwood, W. Ma, F. Moghadam, H.J. Snaith, T. Buonassisi, Z.C. Holman, S.F. Bent, M.D. McGehee, 23.6%-efficient monolithic perovskite/silicon tandem solar cells with improved stability, Nat. Energy. 2 (2017) 17009. https://doi.org/10.1038/nenergy.2017.9.

For high-efficiency CIGS perovskite tandems, Han and coauthors [142] used Cs0.09FA0.77MA0.14Pb(I0.86Br0.14)3 (Eg e 1.60 eV) with monolithical integration, and boosted the original PCE of CIGS subcell from 18.73% to 22.43%, with current density balance at 17.76 mA/cm2 (Fig. 5.24). Shen and coworkers [143] developed a mechanically stacked perovskite/CIGS tandem solar cell, where 1.62 eV band gap perovskite was used with four-cation Cs0.05Rb0.05FA0.765MA0.135PbI2.55Br0.45 composition, with final PCE ¼ 23.9% and increased PCE with respect to the two subcells (PCE ¼ 16.5% for CIGS and 18.0% for PSC) (Fig. 5.24). Concerning the manufacture of tandem solar cells, we can point out that the fabrication processes of the top PSC in 2T tandem should be compatible with the bottom cell. Silicon heterostructure solar cell consists of Si crystals that support temperature up to 500 C and a-Si passivation layers that degrade for T > 200 C. This temperature limitation excludes the high temperature sintering processes needed for transport layers used in PSCs, such as

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Figure 5.24 Device schematic of mechanically stacked perovskite-CIGS tandem with cross-sectional image (A), balanced spectral splitting between CIGS and perovskite subcells (B). CIGS, copper indium gallium diselenide. Reprinted from H. Shen, T. Duong, J. Peng, D. Jacobs, N. Wu, J. Gong, Y. Wu, S.K. Karuturi, X. Fu, K. Weber, X. Xiao, T.P. White, K. Catchpole, Mechanically-stacked perovskite/CIGS tandem solar cells with efficiency of 23.9% and reduced oxygen sensitivity, Energy Environ. Sci. (2018). https://doi.org/10.1039/ c7ee02627g.

TiO2 or NiO. Thus, low-temperature PSC processes are required, which can be obtained by solution processes or vacuum deposition. PSC fabrication requires the deposition of a top transparent electrode to ensure the transmission of infrared light to the bottom subcell. To date, conductive oxides, deposited via magnetron sputtering (MS) processes, are still the best solution for transparent electrodes, in comparison to alternative nanostructured materials such as nanoparticles, carbon nanotubes, metal nanowires, etc. We should point out that high-performing organic transport materials such as Spiro-OMeTAD or PTAA (for HTL) or PCBM and ITIC (for ETL) can be damaged by the high-energy ion flux while depositing TCO with radiofrequency or DC MS methods [156e158]. For this reason, organic transport materials should be completely replaced with an inorganic one, or a special protection barrier should be deposited to organic TL prior electrode deposition. The use of molybdenum oxide (MoOx) as a buffer layer is a standard solution to protect the underlying organic layer [159e161] but requires accurate control of environment conditions such as humidity.

5.4.3 Special structures The introduction of a so versatile perovskite material inspired also alternative architectures which have not been considered in other PV technologies. In the following, we will discuss some of them.

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5.4.3.1 Devices with carbon electrode One of the most crucial interface instabilities in PSCs is the diffusion of metal into the bulk of transport and absorber films that cause shorts, an increased level of nonradiative recombination and formation of metal complexes. In alternative to commonly used back metal electrodes, such as Au, Ag, Al, Cu, conductive carbon was used, that provided, together with an engineered 2D/3D perovskite, a very stable PSC (>10,000 h of stable device operation under light soaking) [64]. In general, device structures with carbon electrodes consider the following stack: FTO/compact TiO2/mesoTiO2/ZrO2 or Al2O3 scaffold/perovskite/carbon [162,163] (see Fig. 5.25). In this configuration, all the layers are printed and then sintered, while the perovskite is infiltrated afterward. Typical thickness for carbon electrodes is in the range of several micrometers (Fig. 5.25B). Solution processing of carbon is not compatible with use of HTL on the top of perovskite, and the absence of electron blocking at anode increases recombination rates and decreases device efficiency, with reduced Voc and high series resistance [152]. Also, the deposition of carbon layers forms poor contact with perovskite film, affecting the whole collection [153]. 5.4.3.2 Resonant NP for light harvesting management Inducing of plasmonic effects is one of the promising efforts for absorption enhancement and photoconversion in solar cells. Liyang et al. [154] (A)

(B)

Figure 5.25 Device schematics of mesoscopic perovskite solar cell (PSC) with ZrO2 scaffold and carbon electrode; (B) Cross section of the ms-PSC with carbon electrode. Reprinted from A. Mei, X. Li, L. Liu, Z. Ku, T. Liu, Y. Rong, M. Xu, M. Hu, J. Chen, Y. Yang, €tzel, H. Han, A hole-conductor-free, fully printable mesoscopic perovskite solar cell M. Gra with high stability, Science (80-.). (2014). https://doi.org/10.1126/science.1254763. (B) Reprinted from K. Lee, J. Kim, H. Yu, J.W. Lee, C.M. Yoon, S.K. Kim, J. Jang, A highly stable and efficient carbon electrode-based perovskite solar cell achieved: via interfacial growth of 2D PEA2PbI4 perovskite, J. Mater. Chem. (2018). https://doi.org/10.1039/c8ta09433k.

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considered a full-wave simulation approach to investigate the contributions of light absorptions across the device illuminated with solar light. Authors revealed that the absorption of the perovskite layer is dominant in UVVis spectral part, while the electrode layers contribute in the red and IR regions. These results demonstrate that it is possible to use silver plasmonic nanoparticles for 58.2% enhancement of the light absorption in the infrared region, by placing a 140-nm-diameter particle arrays (2 nm spacing) via highly localized plasmonic near-fields. Hajjiah and coauthors [155] performed a more detailed modeling for a semitransparent device by placing gold and silver nanoparticles with a diameter of w40 nm at rear electrode and showed a Jsc/EQE enhancement. Furasova and coworkers experimentally realized a novel approach for advanced light harvesting [156]. Si nanoparticles of w50 nm were inserted at the heterojunction interface between perovskite absorber and meso-TiO2 for light trapping in NIR region induced by Mie resonance. The improved device showed 18.8% PCE with >1% of efficiency increase in comparison to the reference cell (Fig. 5.26). Aeineh et al. introduced core/shell Au@SiO2 structures with 10-nm size between compact and meso TiO2 layers, and improved photoconversion. ((A))

( ) (B)

Figure 5.26 Device schematics of mesoscopic perovskite solar cell (PSC) with ZrO2 scaffold and carbon electrode. (B) Cross section of the ms-PSC with carbon electrode. Reprinted from A. Mei, X. Li, L. Liu, Z. Ku, T. Liu, Y. Rong, M. Xu, M. Hu, J. Chen, Y. Yang, €tzel, H. Han, A hole-conductor-free, fully printable mesoscopic perovskite solar cell M. Gra with high stability, Science (80-.). (2014). https://doi.org/10.1126/science.1254763. (B) Reprinted from K. Lee, J. Kim, H. Yu, J.W. Lee, C.M. Yoon, S.K. Kim, J. Jang, A highly stable and efficient carbon electrode-based perovskite solar cell achieved: via interfacial growth of 2D PEA2PbI4 perovskite, J. Mater. Chem. (2018). https://doi.org/10.1039/c8ta09433k.

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Devices with modified ETL with resonant nanoparticles exhibited higher external quantum efficiency in all the visible wavelength range with improved Jsc.

5.5 Stability of PSCs 5.5.1 Light-induced degradation Stability is one of the key challenges for industrial scale commercialization of PSCs. Degradation mechanisms depend on materials, bias conditions, light soaking, and temperature [157]. One of the first studies on metal halide PSCs illustrated that encapsulated MAPbI3 films can be exposed to 1000 h of full sun spectrum, simulated with AM1.5 irradiation, without showing any sign of decomposition as shown by perovskite absorption spectra. This result has been supported by many reports of cells maintaining their full current generation over hundreds of hours of light exposure [151,170,171]. In many studies, however, it is shown that significant changes occur in perovskite films during illumination, such as, ion migration, halide segregation, and compositional degradation [172,173]. While these phenomena present challenges to the perovskite community, photoinduced changes may also be beneficial. For example, in CIGS solar cells, light soaking causes metastable defect reactions with generated charge carriers which lead to increase conductivity, fill factor, and open circuit voltage [174,175]. In addition to the absorber layer, the charge transport layers used in PSCs show variable photostability, resulting in limiting the stability of some types of PSCs. 5.5.1.1 Photostability of charge transport layers Organic semiconductors can form free radicals in the excited state, which can be paired with breaking of carbonenitrogen, carbonecarbon, and carboneoxygen bonds, typically on side chains of the conjugated backbone [175,176]. Under inert conditions, this can result in cross-linking of various free radical species and increases the disorder in the semiconductor films. In fullerene-based materials, which are used as ETLs in PSCs, the photodimerization under photoexcitation creates disorder and inhibits charge transport [164]. Indeed, solar cells using both small molecule hole transport materials and fullerene as electron acceptor materials have been demonstrated to perform stably after hundreds of hours of continuous illumination under 1 sun [151,177,178]. It has been shown that PSCs are robust against minor

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photoinduced degradation within the organic charge transport layers. TiO2 ETLs have been reported to be extremely unstable to UV light not only in the presence of oxygen but even in inert conditions. Photoinduced desorption of oxygen from oxygen vacancies on TiO2 surfaces results in rapid trap-induced recombination across the TiO2 interface [167]. Bella et al. demonstrated that by applying a downconverting fluoropolymer able to transmit the incoming UV light onto the perovskite film as visible light, it is possible to ensure that the TiO2 layer cannot be photoexcited and no photocurrent is lost from the device [168]. 5.5.1.2 Effects on ion distribution in metal halide perovskites DeQuilettes et al. showed that the photoluminescence (PL) intensity of MAPbI3 layers increases over time under 1-sun equivalent irradiation intensity [169]. This increasing of PL intensity is correlated with migration of I species away from the illuminated area. Some of the first direct physical evidence migration of iodide in metal halide perovskite films were provided by this group. This effect could be due to the slow diffusion of ionic species throughout the perovskite layer, and also it is probably related to the presence of point defects, particularly halide vacancies. This group has also found that combining light exposure with oxygen and humidity exposure resulted in even more drastic improvements to PL intensity, causing internal PL quantum yields approaching 90%. Open circuit voltages have been increased in corresponding PSCs [170]. An additional passivating effect of superoxide molecules is suggested, where molecular oxygen can adsorb to surface-trapped electrons and form a stable and passivating superoxidedefect complex by moving the defect energy level outside the band gap. During exposure to humidity and light, an amorphous shell of passivating PbI2 is likely to form, bringing long-lived passivation. 5.5.1.3 Light-induced halide segregation There are many efforts to tune the band gap of PSKs by bromide substitution at halide sites and to increase the thermal stability through Cs and FA substitution on the small cation site. In doing this, however, new lightinduced phenomena have been perceived [66,183]. Under light exposure, Pb(BryI1y)3 perovskites undergo reversible phase segregation into Br- and I-rich perovskite phases; indeed, this observation well aligns with the observation that migration of halide species in perovskite films is induced by light [173]. By removing the light, the materials regain their normal compositional distribution. This phase segregation means that the smaller band gap

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iodine-rich phase inclusions are effectively carrier trapping domains, as confirmed by the fact that PL comes from a low-energy state in such phase-segregated materials. This phenomenon named as the “Hoke effect” provided some definitive evidence that the ions in the perovskite structure can be mobile and that phase segregation can happen in materials having both Br and I on the X site. 5.5.1.4 Light-induced cation segregation There are many studies on cation segregation of the perovskite layer against irradiation [185,186]. Christians et al. revealed that, while the halide and lead distributions remained unchanged all over the depth of the triple catione based PSC, there was a significant shift in Csþ content from the bulk of the material to the hole transporter after 25 h of illumination [174]. The formamidinium and methyl ammonium cations were also redistributed throughout the bulk of the device, which shows further evidence for light-induced compositional changes in PSCs. Large changes in compositional distribution were observed after several hours of operation under full illumination by Domanski et al. [175], with a reversible relative drop in efficiency, up to 15%. The combination of both light and electrical bias was found to be essential to observe the phenomenon. For improvement of the structural stability of the perovskite material, A-site engineering has been applied, but it is critical to specify whether the improved structural stability in the dark translates to structural stability in the light or whether miscibility gaps appear in the compositional spaces for materials in their excited states [176]. 5.5.1.5 Photochemical reactions Metal halides are known to photodecompose. Photodecomposition of silver halide is the most obvious example, which decomposes to halogen and silver upon photoexcitation and formed the basis for early photography [177]. A similar effect has been proposed in PbI2, where carriers trapped at iodide vacancies oxidize iodide to iodine and reduce Pb2þ to Pb0 [178]. Recently, a similar reaction mechanism has been proposed for lead halide perovskites (Fig. 5.27). Kim et al. showed that the halide vacancy concentration dramatically increases by illumination, causing an increased ionic conductivity [161]. This observation is consistent with other studies reporting a rising in ionic conductivity under illumination [179]. They suggest that iodide ions in the lattice are oxidized by photogenerated holes, resulting in coupled formation of neutral iodine interstitials and iodide vacancies.

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Figure 5.27 (A) Photodecomposition mechanism. (B) The movement of ions and the formation of molecular iodine and metallic lead upon illumination. (A) Reprinted from M.G. Albrecht, M. Green, The kinetics of the photolysis of thin films of lead iodide, J. Phys. Chem. Solids 38 (1977) 297e306. https://doi.org/10.1016/0022-3697(77)90106-8. (B) €m, V. Lanzilotto, F.O.L. Johansson, K. Aitola, B. Reprinted from U.B. Cappel, S. Svanstro €hlisch, S. Svensson, N. MårPhilippe, E. Giangrisostomi, R. Ovsyannikov, T. Leitner, A. Fo tensson, G. Boschloo, A. Lindblad, H. Rensmo, Partially reversible photoinduced chemical changes in a mixed-ion perovskite material for solar cells, ACS Appl. Mater. Interfaces. 9 (2017) 34970e34978. https://doi.org/10.1021/acsami.7b10643.

5.5.2 Reactions with electrodes Au, as electrode in PSKs devices, does not form a redox couple with perovskite while some metals can form redox couples with the perovskite itself, even reacting with PbI2 [180]. Some of them erode in the presence of reactive polyiodide melts, formed from perovskite decomposition in visible light [179,192]. Thus, while some metals may be stable with respect to the perovskite structure, almost all metals react with decomposition products of the perovskite, such as HI, MAI, CH3I, and I2, that could be produced even in an encapsulated PSC. For metal contacteinduced degradation in PSCs, there are three major mechanisms, all of which cause the device performance to drop significantly: (1) halide anions diffuse to the metal electrode, corroding the metal and resulting in a halide deficiency in the perovskite absorber layer (Fig. 5.28A) [182]; (2) metal diffuses under activation of light and/or heat into the perovskite film, potentially forming insulating metal halide species or defect states in the bulk or at the perovskite interface (Fig. 5.28, panels b and c) [183]; (3) metal contacts form a redox couple with Pb2þ in perovskite layer, accelerating the loss of halide species and forming Pb0 [180].

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Figure 5.28 (A) Schematic detailing reaction of a metal electrode with halide/halogen species created via decomposition of the perovskite by moisture. (B) TOF-SIMS depth profile showing diffusion of Au into the perovskite film after heating at 70  C under illumination, with (C) a corresponding 3D elemental map. (A) The figure has elected from Y. Kato, L.K. Ono, M. V. Lee, S. Wang, S.R. Raga, Y. Qi, Silver iodide formation in methyl ammonium lead iodide perovskite solar cells with silver top electrodes, Adv. Mater. Interfaces. 2 (2015) 1500195. https://doi.org/10.1002/admi.201500195. (B) Reprinted from K. Domanski, J.-P. Correa-Baena, N. Mine, M.K. Nazeeruddin, A. Abate, M. Saliba, W. Tress, A. €tzel, Not all that glitters is gold: metal-migration-induced degradation in Hagfeldt, M. Gra perovskite solar cells, ACS Nano. 10 (2016) 6306e6314. https://doi.org/10.1021/acsnano. 6b02613.

5.6 Upscaling of perovskite solar devices The devices of the second [192e198] and third [199e222] generations of PV are based on transparent conducting oxides (TCOs) [223], employed as photoelectrodes. These materials are characterized by a good degree of transparency and relatively low sheet resistances, typically in the order of 7e15 U/, [224e226]. In the case of large-area devices, the latter values could be deleterious for the electrical performance, generating a high series resistance that opposes to current flux. Therefore, to realize large-area PSMs, it is mandatory to divide full-sized TCO substrates into multiple cells. Two main architectures could be employed, one that uses series-connected cells and the other that uses parallel-connected cells.

5.6.1 Series-connected solar modules The architecture that is principally employed to fabricate PSMs is the seriesconnected one [227e236]. In this case, the voltage of the cells will be ideally

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Figure 5.29 Cross section of a typical thin-film series-connected solar module. The P1, P2, and P3, discussed below, constitute a zone called dead area ([, in the figure), dedicated to the interconnection between subsequent cells. The photoactive area is here indicated with L. RS, top and RS, bottom represent the series resistance offered by the two electrodes, Rint represents the contact resistance between the electrodes of subsequent cells in the series architecture. Reprinted from L. Lucera, F. Machui, P. Kubis, H.J. Egelhaaf, C.J. Brabec, Highly efficient, large area, roll coated flexible and rigid solar modules: design rules and realization, in: Conference Record of the IEEE Photovoltaic Specialists Conference, 2016, pp. 234e237, https://doi.org/10.1109/PVSC.2016.7749585. Copyright (2016) IEEE.).

summed to obtain the voltage of the module, while the current of the latter will be dominated by one of the less performing cells in the series. In Fig. 5.29, different resistances are shown, two generated by the electrodes and one that characterizes the contact between the TCO of a cell and the counter-electrode of the subsequent one.

5.6.2 Parallel-connected solar modules In order to raise the current output of the module, while, ideally, maintaining the same voltage of a single cell, the parallel-connection architecture can be employed. The latter is normally realized maintaining the same TCO layer for all of the cells of the module. Due to the TCO resistivity, above introduced, this leads to a strong reduction of performance [237e239], which has brought to a scarcer scientific investigation of the parallel-connection architecture with respect to the series-connection one [237,240,241]. Different stratagems have been adopted to improve the performance of parallel-connection architecture, e.g., the employment of a metal grid printed over the TCO, to increase the charge collection on the photoelectrode [240,241], as shown in Figs. 5.30 and 5.31.

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(A)

(B)

Figure 5.30 Device structure (A) and layout (B) by Hambsch et al. [240]. The aluminum grid was deposited directly onto the ITO layer by thermal evaporation and, then, exposed to an UV-ozone plasma to realize an aluminum oxide layer; the latter improves the hydrophilic property of the film, reducing the possibility of device short circuit between the grids and top electrode and can prevent corrosion of aluminum grids from the perovskite film. Adapted from M. Hambsch, Q. Lin, A. Armin, P.L. Burn, P. Meredith, Efficient, monolithic large area organohalide perovskite solar cells, J. Mater. Chem., 4 (2016) 13830-13836, https://doi.org/10.1039/C6TA04973G. Copyright (2016) The Royal Society of Chemistry.

Figure 5.31 Device layout and certified IeV curve by Kim et al. [241] The authors did not specify which metal has been employed for the grid, but only its resistivity ¼ 9.3$106 U cm. Adapted from J. Kim, J.S. Yun, Y. Cho, D.S. Lee, B. Wilkinson, A.M. Soufiani, X. Deng, J. Zheng, A. Shi, S. Lim, S. Chen, Z. Hameiri, M. Zhang, C.F.J. Lau, S. Huang, M.A. Green, A.W.Y. Ho-Baillie, Overcoming the challenges of large-area highefficiency perovskite solar cells, ACS Energy Letters, 2 (2017) 1978e1984, https://doi.org/ 10.1021/acsenergylett.7b00573. Copyright (2017) American Chemical Society.

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The optimization of such a layout has brought to a maximum certified PCE ¼ 12.1% on a device with AA ¼ 16 cm2 [241]. This result sets a record for PSMs based on parallel-connected architecture. Nevertheless, the series-connected architecture evidently makes the scalability of devices easier, and a wider research activity has been focused on series-connected architecture, which holds all of the current records for each dimension-category of PSMs [228,235,242]. To realize the interconnections of such an architecture, the P1eP2eP3 process is employed.

5.6.3 The P1eP2eP3 process In a series-connected PSM, due to the considerable sheet resistance of TCO, a subdivision of the substrate into subcells is necessary, as previously mentioned. In this way a monolithic interconnection scheme [228] is realized, as shown in Fig. 5.32 [236]. The latter is constituted by some active zones that are dedicated to the photogeneration of electric charges and, thus, to the PV energy conversion, and dead zones, dedicated to the interconnection between adjacent cells. It is possible to define the active area (AA) as that part of the substrate dedicated to the PV energy conversion and the dead area (DA) as the part where the interconnection of cells takes

Figure 5.32 Schematic representation of a typical mesostructured perovskite solar module. In the zoomed circle, an active area width (WA) and a dead area width (WD) are highlighted. The latter is formed by the P1, P2, and P3 processes lines and by the two safety areas (SA) between these. Reprinted from A.L. Palma, F. Matteocci, A. , L. Vesce, S. Christiansen, M. Schmidt, A.D. Carlo, LaserAgresti, S. Pescetelli, E. Calabro patterning engineering for perovskite solar modules with 95% aperture ratio, IEEE Journal of Photovoltaics, 7 (2017) 1674e1680, https://doi.org/10.1109/JPHOTOV.2017. 2732223. Copyright (2017) IEEE.

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place, and which does not contribute to the energy conversion. The DA is formed by the P1, P2, and P3 zones and by the two safety areas (SA) that are interposed between these. The respective functions will be described in the following paragraphs. In Fig. 5.32 [236], the relative widths (WA, WD ¼ P1þSA þ P2þSA þ P3) are highlighted. It is now possible to define the aperture ratio (AR) as AR ¼

WA WD þ WA

(5.2)

where the aperture area width is defined as the sum of WD and WA. It has been shown that the interconnections in series-connected PSMs can be the realized employing several patterning approaches [228,231,233,234]. Nevertheless, laser processing has been demonstrated to be an optimal choice for industrial applications, already for secondgeneration technologies [221,243e246], since it provides for the best trade-off between manufacturability and minimization of DAs [234]. The P1 process insulates the photoanodes of neighboring cells, for series connection. The P2 patterning step selectively removes the entire stack of active materials between adjacent cells to permit their series connection through the subsequent deposition of a conductive layer. Finally, the latter is homogeneously deposited on the entire module and the P3 step is applied to separate the just realized adjacent counter-electrodes. The P1eP2eP3 process on PSMs has to be performed taking into account the geometrical losses due to the presence of DAs, the resistive losses presented by the discrete sheet resistance of the TCO, and the ohmic contact losses due to the limited width of the interconnection areas and to the characteristic transfer length (LT) [228,247e250] of the realized contact between the electrodes. These can be resumed in a parameter, named total loss (TL) [251,252]. The TL can be expressed as a function of AA width (WA), DA width (WD), LT, and, in particular, P2 process width. By extending the procedure in Ref. [252] and according to well-known mathematical models [251,253,254], the TL has recently been defined, in terms of the modules parameters [236], as: Total Loss ¼

TCO R, WD J WA3 þ MPP WA þ WD VMPP 3 WA þ WD  ! WA2 P2 P2 þ LT R, coth LT WA þ WD

(5.3)

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where JMPP and VMPP are the current density and voltage at maximum TCO is the sheet resistance of the power point of a small area single cell, R, P2 active area TCO photoelectrode, R, is the sheet resistance of the TCO in the interconnection area after the realization of the P2 process, which could possibly modify the TCO surface. The realization of a P1eP2eP3 procedure is represented in Fig. 5.33. 5.6.3.1 P1 process The P1 is the first process in the fabrication of PSMs. It is realized to create the photoelectrodes of the cells, starting from a substrate homogeneously covered by a TCO. In PSMs, the P1 is typically realized with the application of IR laser machines that achieve the ablation through thermal vaporization of the FTO [228e230,235,236,255] or ITO [234,242] on glass substrates and ITO on plastics substrates [256,257].

Figure 5.33 Production process flow of a typical p-i-n structured perovskite solar module based on compact and mesoscopic TiO2 ETM.

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The width of P1 ablation lines should be the thinnest that still realizes an isolation between adjacent TCO zones. In the fabrication of PSMs, P1 widths in the range of 40e50 mm have been reported [236,258]. 5.6.3.2 P2 process The most delicate step in the realization of PSMs interconnections is the P2. The latter is actuated after the deposition of all of the active materials on the TCO and just before the deposition of the counter-electrode. The P2 step is needed to remove the active stack from the zones of TCO dedicated to the interconnection with subsequent cells in the series and that have to be preserved intact. Apart from very few mechanical microblade scribing techniques [233,234], which are complicated and not scalable, laser processing has been the most employed technique for the realization of the P2 step. In fact, laser techniques are wavelength selective and permit to ablate the active materials, which are highly absorptive in the visible range of the light spectrum, from the TCOs, which, instead, are lowly absorptive in the same wavelengths range. Therefore, the mainly employed laser wavelengths for the P2 step have been in the range of 515e532 nm [244,259] and 355 nm [236,242]. Different kinds of sources have been used, typically in the nanosecond (ns)-range [246], and picosecond (ps)-range [236,242], like fiber lasers [260], and diode-pumped solid-state Q-switched lasers [244]. Due to the discrete resistivity of the TCOs, the P2 scribe width cannot be indiscriminately reduced, like in the P1 case, to avoid high contact resistances between the electrodes. For PSMs, P2 dimensions in the range of 45e200 mm have been reported [231,234,236,242]. 5.6.3.3 P3 process After the P2 step, a homogeneous counter-electrode material layer is deposited and, then, the P3 is realized to isolate, one to each other, the counterelectrodes of the single cells. Typically, the same laser systems and parameters employed for the P2 step can also be used in the P3 step realization [234,236]. As for the P1, the P3 scribing can be performed realizing the thinnest width that still isolates the electrodes. In PSMs fabrication, P3 process widths as low as 25 mm have been reported [236]. 5.6.3.4 Safety areas While realizing a P1eP2eP3 procedure, in the steps between the laser processes, when the substrate is removed from the laser stage to deposit the

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materials, some possible misalignments are introduced, with respect to the position assumed during the P1 step. Therefore, even considering top-inclass subsequent alignment steps, the realization of SAs between P1eP2 and P2eP3 is necessary to avoid detrimental overlapping of these steps (see Fig. 5.32). SAs as thin as 50 mm have been reported [45].

5.6.4 Deposition techniques The morphology of the perovskite layer plays a crucial role in order to obtain high-efficiency PSCs. Several deposition techniques have been used for the perovskite deposition such as spin coating, blade coating, slot-die coating. The main constraints are related to the optimization of the solutions and deposition parameters in order to obtain highly uniform perovskite layer in terms of growth and substrate coverage. All the layers of a perovskite-based device can be deposited starting from solutions. Two main deposition strategies can be adopted, the sheet-to-sheet production and the roll-to-roll production. The first one involves a single glass or plastic substrate, on which the materials are deposited. The most used deposition techniques are the spin coating, the blade coating, or the slot-die coating. The spin coating is mostly used for small-area solar cells and modules rapid prototyping since it shows high thickness and film quality control, obtained by setting a series of parameters like the acceleration time, speed, and time of rotation. Nevertheless, while increasing the dimensions of the substrates, this technique has been shown not to be optimal for an industrial-oriented production, mostly for the impressive rate of material that is wasted during such a deposition process. An alternative way of depositing active layers of a large-area PSM is the blade coating (or doctor blade). This technique can be adopted both on rigid and on flexible substrates. In the sheet-to-sheet processes, typically, the substrate is fixed and a liquid solution is deposited on an external part of the latter; then, a blade moves over it, spreading the liquid precursor [229,261]. The precursor solution, the speed of the blade, and the height of the latter from the substrate are parameters that can be set. The slot die coating is one of the most promising deposition techniques. In this case, a liquid solution passes through the slot-die head that, moving over the substrate, distributes the precursor on it. Respect to the blade coating, the distribution and the deposition of the solution are part of the same process.

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Figure 5.34 Schematics of a roll-to-roll production system that includes a slot-die coating process. Reprinted from B. Dou, J.B. Whitaker, K. Bruening, D.T. Moore, L.M. Wheeler, J. Ryter, N.J. Breslin, J.J. Berry, S.M. Garner, F.S. Barnes, S.E. Shaheen, C.J. Tassone, K. Zhu, M.F.A.M. van Hest, Roll-to-Roll printing of perovskite solar cells, ACS Energy Letters, 3 (2018) 2558e2565, https://doi.org/10.1021/acsenergylett.8b01556. Copyright © 2019 American Chemical Society.)

The roll-to-roll production takes inspiration from the journals printing processes. As depicted in Fig. 5.34 [262], this production procedure involves rolls that move a flexible substrate along a path. Distributed on the latter, different production devices are placed, such as the deposition part, like a blade coating or a slot-die machine, the zone dedicated to the annealing, such as a double-sided furnace; a zone can be dedicated to the laser patterning. The roll-to-toll production procedure, clearly, permits to reach a higher throughput respect to the sheet-to-sheet one and has been investigated to reach an industrially oriented production rate for perovskite-based largearea modules [263e265].

5.7 Conclusions and perspectives This chapter provided a brief introduction to working principles and main directions for the development of PSC technology. The unprecedented rapid progress in the research of halide perovskites for PV

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applications during the last decade demonstrates real industrial potential for a cost-effective, low capital expenditure (CAPEX) printed thin-film technology that can be the competitive to traditional semiconductors such as Si, CIGS, CdTe, etc. At the same time, the unique semiconductor characteristics, related to band structure properties, bipolar charge transport, ionic motion, and traps still need to be clarified and related to the exceptional performance of these materials. The main strategies for future development of PSC are oriented to the optimization of photo and thermal stability of absorber and transport layers. In addition, interface engineering for improvement of charge extraction, passivation of perovskite grain-boundaries/surfaces, and to prevent metal electrode diffusion needs to be further developed to achieve an international standard assessment of PSC stability. In turn, this will open industrial exploitation of perovskite PV technology. In this context, the possibility to tune on demand the optoelectronic properties of the absorber, by varying the halide perovskite composition, opens promising perspectives for spectral management in tandem integration with silicon and CIGS PVs. The PCE of 28% achieved so far for perovskite/silicon tandem allows us to forecast a bright future of the emerging technology of perovskite multijunction solar cells.

Acknowledgments D.S. and A.D.C. gratefully acknowledge the financial support of the Ministry of Science and Higher Education of the Russian Federation in the framework of Increase Competitiveness Program of NUST «MISiS» (N L2-2019-013), implemented by a governmental decree dated 16th of March 2013, N 211.

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[263] Y. Galagan, F. Di Giacomo, H. Gorter, G. Kirchner, I. de Vries, R. Andriessen, P. Groen, Roll-to-Roll slot die coated perovskite for efficient flexible solar cells, Advanced Energy Materials 8 (2018) 1801935, https://doi.org/10.1002/ aenm.201801935. [264] D. Burkitt, J. Searle, A.D. Worsley, T. Watson, Sequential slot-die deposition of perovskite solar cells using dimethylsulfoxide lead iodide ink, Materials (2018) 11, https://doi.org/10.3390/ma11112106. [265] F. Di Giacomo, H. Fledderus, H. Gorter, G. Kirchner, I. de Vries, I. Dogan, W. Verhees, V. Zardetto, M. Najafi, D. Zhang, H. Lifka, Y. Galagan, T. Aernouts, S. Veenstra, P. Groen, R. Andriesse, Large area >140 cm^2 perovskite solar modules made by sheet to sheet and roll to roll fabrication with 14.5% efficiency, in: 2018 IEEE 7th World Conference on Photovoltaic Energy Conversion (WCPEC) (A Joint Conference of 45th IEEE PVSC, 28th PVSEC & 34th EU PVSEC), 2018, pp. 2795e2798, https://doi.org/10.1109/PVSC.2018.8548157. [266] L. Lucera, F. Machui, P. Kubis, H.J. Egelhaaf, C.J. Brabec, Highly efficient, large area, roll coated flexible and rigid solar modules: design rules and realization, in: Conference Record of the IEEE Photovoltaic Specialists Conference, 2016, pp. 234e237, https://doi.org/10.1109/PVSC.2016.7749585.

CHAPTER SIX

All-oxide solar cells Mojtaba Gilzad Kohan1, Isabella Concina1, Alberto Vomiero1, 2 1

Division of Materials Science, Department of Engineering Sciences and Mathematics, Luleå, Sweden Department of Molecular Sciences and Nanosystems, Ca’ Foscari University of Venice, Italy

2

6.1 Introduction In the quest for the development of the next-generation solar cells, the research aims at addressing simultaneously three main paradigms, relying on the possibility of obtaining (1) high efficiency devices made of (2) environmentally friendly and low-cost materials exhibiting (3) long-term stability of their functionalities. Different technologies exist, fulfilling separately one or two of the aforementioned features. For instance, high-efficiency multijunction solar cells exhibit maximum photoconversion efficiency (PCE) around 47% [1], but their use is limited to niche applications due to their high cost. Perovskite solar cells [2], aside of having high PCE (record certified efficiency close to 23%) and low cost, still suffer for their poor stability and the use of environmentally harmful materials. Organic photovoltaics exhibit good photoconversion efficiency (max PCE around 15.6% [1]) and low cost but are inherently unstable, limiting their use to short-term applications and devices. One of the original directions in this panorama addresses the study of solar cells based on oxide materials, which can be inherently stable and environmentally friendly, and are typically prepared by scalable and low-cost synthetic routes. Metal oxides are a broad class of materials, in which both intrinsically p- and n-doped semiconductors can be found. The electronic properties in metal oxides allow the existence of wide band gap semiconductors (like, for instance, ZnO, TiO2, and SnO2) with energy gaps Egw3.2 (ZnO, TiO2)e3.8 eV (SnO2), which are transparent to visible solar radiation and small band gap semiconductors (like, for instance, Co3O4, CuO, and Cu2O) with energy gaps Egw(0.7, 1.5, 2.1, and 4.4) eV (Co3O4) [3], 2.1 eV (CuO) [4], 1.2 eV (Cu2O) [4], able to absorb solar light in the visible and near-infrared (NIR) spectral region. Interestingly, Egw1.5 eV for Co3O4 [5] is very close to the ideal Egw1.34 eV for the thermodynamic limit of PCE according to the Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00006-9

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ShockleyeQueisser calculation [6,7], which suggests this oxide to be very promising to obtain highly efficient devices [8]. Another good property of most of the oxides is their high extinction coefficient [9]. This is a critical feature for the development of solar cells, since it requires a limited thickness of the optically active layer to get full absorbance of sunlight radiation. Limited carrier mobility (especially in p-type oxides) induces strong charge recombination, so the possibility of having thin absorbers is a benefit. For all these reasons, existence of optimum Eg, scalability of preparation processes, eco-friendliness and low cost of the proposed materials, all-oxide solar cells are an interesting class of next-generation solar cells, which are attracting significant interest in the last years [10,11]. The main configuration in which these materials are studied is the pen single heterojunction device: the built-up potential of the pen junction is used to separate the exciton created by light absorption, similarly to what happens in a pen homojunction silicon cell. Unfortunately, up to now, rather limited PCE was demonstrated for this class of devices, which slightly exceed the 6% [12] (see Table 6.1 from Ref. [11] for a comparative analysis of the various photoconversion efficiencies in all-oxide solar cells).

6.2 Electronic band structure, interface tuning, and doping In the process of light photoconversion, three main distinct physical mechanisms can be identified, namely light absorption and exciton creation, exciton separation at the pen junction, charge transport and collection at the external electrodes. The first requirement for an operating device is to have a suitable electronic band structure in the junction, which makes it possible to separate the exciton created during photon absorption, having the electron injected into the n-type layer and the hole transported through the p-type semiconductor. An example of suitable electronic structure for a pen junction is reported in Fig. 6.1, from Ref. [24]. In their work, Atwater and coworkers demonstrated the critical role of conduction and valence band positions in determining the photovoltaic features of pen heterojunctions. Indeed, while in homojunction pen silicon devices the electronic band alignment between the p- and n-regions is trivial, it is not the case in heterojunction devices, where unwanted band misalignment may fully inhibit the operational properties of the device.

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Table 6.1 Functional properties of oxide solar cells based on different composite materials and geometries. Jsc Fill factor PCE Voc (mV) (mA cmL2) (%) (%) References System

Mg2/AZO/ (Ga0.975Al0.025)2O3 /Cu2O:Na Al:ZnO/Ga2O3/Cu2O ZnO:Al/a-ZTO/Cu2O ZnO/a-ZTO/Cu2O Zn0.79Mg0.21O/Cu2O Cu2O/Cu p-CuO:Al/n-Si NiO/Cu2O/ZnO/ SnO2 Cu2O homojunction ZnO NRs (700 nm)/ Cu2O Cu2O/ZnO NWs Cu2O/ZnO NWs ZnO/Cu2O ZnO/Cu/Cu2O Co3O4/TiO2

840

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800 622 553 650 595 340 640

9.99 7.25 7.37 6.9 6.78 8.7 3.5

67 63.1 65.0 49.2 50 68 50

5.38 2.85 2.65 2.2 2 2 1.12

[13] [14] [15] [16] [17] [18] [19]

287 230

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25 33

0.89 0.43

[20] [21]

130 150 530 347 430

2.35 1.43 0.69 1.41 w0.2

29 25 30.4 26.0 /

0.1 0.053 / / /

[22] [23] [24] [24] [8]

Reproduced with permission from I. Concina, Z.H. Ibupoto, A. Vomiero, Semiconducting metal oxide nanostructures for water splitting and photovoltaics, Adv. En. Mater., 7 (2017) 1700706. Copyright 2017 Wiley.

Figure 6.1 Electronic band structure and valence band offset measurements derived from X-ray photoelectron spectroscopy of various films. Reproduced with permission from S.S. Wilson, J.P. Bosco, Y. Tolstova, D.O. Scanlon, G.W. Watson, H.A. Atwater, Interface stoichiometry control to improve device voltage and modify band alignment in ZnO/Cu2O heterojunction solar cells, Energy Environ. Sci., 7 (2014) 3606e3610. Copyright 2014 Royal Society of Chemistry.

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For this reason, suitable choice of the components is mandatory. A clear example for this is demonstrated in Ref. [25], in which an n-type Sn1xCoxO2 and a p-type Cu2O layer are used to form the pen junction. The suitable electronic structure of the junction resulted in a record Voc (2.325 V) under 1 sun illumination (Fig. 6.2), which is among the highest Vocs for any type of solar cell. It was also demonstrated that the interface properties of the heterojunction are very sensitive to the presence of defects and/or spurious phases, and that even the presence of contaminants may degrade the functionality. It was the case, in Ref. [24], of the Cu2O/ZnO thin-film system. The presence of CuO phase at the interface induces a shift in the valence band offset for the ZnO/Cu2O junction, which inhibits proper charge injection from the photoexcited CuO to the ZnO conduction band, impairing the photovoltaic properties under sunlight illumination. As a consequence, strong charge recombination occurs, according to the electronic band structure reported in Fig. 6.1, since a potential barrier inhibits CuO to ZnO electron injection.

Figure 6.2 (A) UV-Vis absorption spectra, (B) Voc, Jsc, (C) PCE, and fill factor (FF) for Sn1xCoxO2/Cu2O solar cells for two different values of Co doping. (D) The JeV curve relates to sample Sn0.86Co0.14O2/Cu2O exhibiting the highest PCE (1.20%). Reproduced with permission from H.Y. Shiu, C.M. Tsai, S.Y. Chen, T.R. Yew, Solution-processed all-oxide nanostructures for heterojunction solar cells, J. Mater. Chem., 21 (2011) 17646e17650. Copyright 2011 RSC.

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From this point of view, the most investigated system has been the one based on the Cu2OeZnO pen heterojunction thin film [26]. Several different synthetic routes were applied for the purpose, from physical vapor deposition (mainly sputtering) to chemical and electrochemical deposition. The suitability of all-oxide solar cells for market applications grounds on the need to produce oxide solar cells through vacuum-free technologies. Originally, a series of studies focused on preparation of oxide solar cells by using a copper oxide layer grown through thermal oxidation of a Cu metallic substrate [16,27]. However, this method has the drawback of large Cu consumption, impairing its practical application. Moving to thin-film technology, avoiding Cu foil substrate, ZnO/ Cu2O heterojunction solar cells were produced through sequential cathodic electrodeposition of ZnO and Cu2O on conductive glass substrates [28]. Rather thick layers were used (w2-mm-thick ZnO and w2-mm-thick Cu2O films) for a total thickness of the cell of about 4 mm. A simple and scalable methodology (spray pyrolysis) was applied for the deposition of a TiO2/Cu2O heterojunction thin-film solar cell [26]. This report represents the first example for a complete cell to be entirely fabricated by easily scalable low-cost technique. The quest for cheap and environmentally friendly methodologies is still under development, and a very recent publication focused on the preparation of all-solution-processed all-oxide solar cell using exclusively green and low-cost raw materials deposited from aqueous solutions [29]. The interesting point of the work is the fine-tuning of the interfacial properties of the p- and n-type layers, which influence the main functional parameters (Jsc, Voc, and FF) of the solar cell. Another direction toward the cost reduction is the application of atmospheric conditions [19] for deposition of Cu2O and CuO (after annealing of the Cu2O film) through successive ionic layer absorption and reaction. As already emphasized, the electronic band structure is highly important. In this work, the electronic band structure was measured directly through scanning tunneling spectroscopy to locate the band edge position of the different constituents with respect to their Fermi energy. The measured electronic band structure is reported in Fig. 6.3. It is experimentally proven that different phases (CuO/Cu2O) modulate the electronic band alignment of the multilayered film, resulting in different functionalities. The clearest signal is a simultaneous and drastic decrease of both short circuit current density and open circuit photovoltage moving from Cu2O to CuO (Fig. 6.3). This work also highlights the importance of the presence of interface and contact layers, which can act either as charge collectors (NiO thin film here) or as selective

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Figure 6.3 Schematic energy level diagram of (A) NiO/Cu2O/ZnO/SnO2 and (B) NiO/ CuO/ZnO/SnO2 heterojunctions. The dashed line represents the Fermi energy after contact. Currentevoltage characteristics of (C) NiO/Cu2O/ZnO/SnO2 and (D) NiO/ CuO/ZnO/SnO2 heterojunctions under a dark and a white light illumination condition. Reproduced with permission from S. Chatterjee, S.K. Saha, A.J. Pal, Formation of all-oxide solar cells in atmospheric condition based on Cu2O thin-films grown through SILAR technique, Sol. Energy Mater. Sol. Cells, 147 (2016) 17e26. Copyright 2016 Elsevier.

blocking layers for either holes or electrons (ZnO thin film here to block hole diffusion from the optically active layer to the SnO2 film). Further improvement of solar cell performance can be obtained by suitable doping of the absorber material [30]. Fortunato and coworkers proposed to dope a Cu2O thin film with Mg. They demonstrated that Mg doping does not influence significantly the optical properties (and specifically the energy band gap) of the Cu2O layer, due to similar ionic radii of Mg2þ and Cuþ, but it affects the crystallite size by probably increasing the number of nucleation centers. This way, formation of phases different from Cu2O is prevented. Mg doping significantly increases the photoconductivity of the Cu2O layer and prevents the formation of trap states, which is one of the main causes of degradation of PCE through limitation of minority carrier diffusion in Cu2O.

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Calcium was another suitable candidate for doping Cu2O thin films [31]. Ca-doping increases the crystalline size and roughness of the p-type Ca: Cu2O polycrystalline film, and lowers the energy gap (from 2.40 to 2.15 eV). The highest PCE (0.45%) was achieved for 3% atomic Ca doping. The critical importance of the quality of the interface was discussed by MacManus-Driscoll and coworkers [16]. They produced ZnO/Cu2O heterojunction solar cells under atmospheric conditions through atmospheric atomic layer deposition of a Zn1-xMgxO film on thermally oxidized cuprous oxide. The formation of spurious CuO phase atop of the Cu2O layer was proven by photothermal deflection spectroscopy (Fig. 6.4), indicating that it can be effectively removed through ion etching in an oxygen-free environment. Removal of the CuO layer significantly increases the PCE up to 2.2%, which is one of the highest PCE for vacuum-free all-oxide solar cells. Nitrogen doping was also demonstrated to be beneficial in terms of hole density increase [32]. The doping was obtained during the in situ growth of Cu2O by sputtering. Hole density was found to increase almost linearly with N2 flow in the sputtering chamber, and a significant increase of two orders of magnitude was observed (from 1015 to 1017 cm3). On the other hand, authors remarked that nitrogen flow rate affects the hole mobility (decreasing with increasing flow rates) within the Cu2O layer (thickness about 2 mm), so 1

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Figure 6.4 Photothermal deflection spectra of Cu2O films before (red) and after (black) etching to remove the CuO surface layer formed due to surface oxidation during thin film growth. The comparison of the two spectra highlights the presence of absorption from CuO in the film before etching. Reproduced with permission from Y. Ievskaya, R.L.Z. Hoye, A. Sadhanala, K.P. Musselman, J.L. MacManus-Driscoll, Fabrication of ZnO/Cu2O heterojunctions in atmospheric conditions: improved interface quality and solar cell performance, Sol. Energy Mater. Sol. Cells, 135 (2015) 43e48. Copyright 2015 Elsevier.

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that nitrogen can be used to modulate the electrical properties of the semiconductor. The nitrogen is thought to occupy oxygen crystal sites and thus acts as an acceptor. In the same investigation, the authors subjected the Cu2O layer to a crown-ether cyanide treatment, which also resulted in increasing the hole density by an order of magnitude. This finding was ascribed to the passivation of nonradiative recombination centers of Cu2O grains through the elimination of Cu dangling bonds on the metal oxide surface. The functional parameters of a working device was found to be 0.26 V as for the VOC, 2.8 mA/cm2 as for JSC, and 0.55 as for the FF, with an overall conversion efficiency of 0.4%. A homojunction n-Cu2O/p-Cu2O was also proposed [33]. In their investigation Wei and coworkers used an electrochemical deposition synthetic strategy that revealed capable of tuning the growth orientation of the nCu2O from to so that the subsequent epitaxial growth of the p-Cu2O may be in the same crystallographic orientation, thus allowing to obtain a high-quality homojunction with reduced interface states (Fig. 6.5). Indeed, the devices exploiting the materials with the same crystallographic orientation were found to show the best performances of the studied batch, although the solar energy conversion was overall quite low (VOC ¼ 0.12 V, JSC ¼ 3.97 mA/cm2, FF ¼ 22.6, PCE ¼ 0.104%).

6.3 Nanostructured architectures and nanowires The trade-off between light absorption and charge collection is particularly severe in oxide solar cells due to the low mobility of minority carriers in most of the p-type oxides. This limits significantly the maximum thickness of the light absorption layer, resulting in severe light loss. In case thicker layers are applied, instead, strong charge recombination occurs, due to poor transport properties. For instance, it has been demonstrated that in Cu2O the inverse of the absorption coefficient is larger than the minority carrier diffusion length, inducing strong charge recombination within the light absorber. To overcome this issue, nanostructured architectures can be beneficial, decoupling the processes of light absorption and charge transport [34]. Nanowire geometries can be usefully applied for the purpose [23]. Nanowires have been applied in the past to boost electronic transport in some kinds of excitonic solar cells, like dye-sensitized [35,36] and quantum dotesensitized solar cells [37]. The main claim in that kind of systems was that the electron diffusivity in ZnO single crystal nanowires was 100 times larger than diffusivity for TiO2 or ZnO mesoporous films, enabling fast and direct electron collection,

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Figure 6.5 (A) Structural scheme of the Cu2O homojunction solar cell. (B) Side-view SEM image of the Cu2O homojunction. (C) Dark IeV characteristics of the Cu2O homojunction. (D) Currentvoltage (IeV) characteristic curves of the Cu2O homojunction solar cells obtained by different deposition potentials. Reprinted from H.M. Wei, H.B. Gong, L. Chen, M. Zi, B.Q. Cao, Photovoltaic efficiency enhancement of Cu2O solar cells achieved by controlling homojunction orientation and surface microstructure, J. Phys. Chem. C, 116 (2012) 10510e10515. with ACS permission. Copyright 2012 ACS.

overcoming the hopping process, which regulates electronic motion in mesoporous nanoparticle films [36]. However, both dye- and quantum dote sensitized cells are based on the absorption of a single layer of the optically active material at the surface of the oxide scaffold. Multiple layer deposition results in strong recombination due to reduced exciton separation. For this reason, the optical density of the electrode is proportional to its specific surface area, which is rather limited in a nanowire array, compared to a mesoporous film. As a consequence, most of the nanowire-based solar cells sensitized through dye or quantum dots suffer from low optical density. This issue can be overcome by implementing a mixed nanowire/mesoporous film network (Fig. 6.6) [38], able to offer the high specific surface of the mesoporous film and the good transport properties typical of the ZnO nanowire array. The optical characterization demonstrated the enhanced light

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Figure 6.6 Nanowire-based photoanodes in excitonic solar cells: (A) Bare nanowire array. (B) Composite photoanode, in which a mesoporous TiO2 nanoparticle layer permeates the ZnO nanowire array. (C) Absorbance and (D) current densityevoltage curves under 1 sun irradiation (AM 1.5G) of the photoanodes sensitized by N719 commercial dye. Solid line: nanowire array; dashed line: composite system. Reproduced with permission from A. Vomiero, I. Concina, M. Natile, E. Comini, G. Faglia, M. Ferroni, I. Kholmanov, G. Sberveglieri, ZnO/TiO2 nanonetwork as efficient photoanode in excitonic solar cells, Appl. Phys. Lett., 95 (2009) 193104. Copyright 2005 American Institute of Physics.

absorption directly related to increased specific surface area in the composite system, compared to the pure ZnO nanowire array, which directly reflected in the increased short circuit current density, from 1.6 to 6.8 mA cm2. While the optical density is an issue in dye- and quantum dotesensitized solar cells, due to the required single-layer deposition, in all-solid state oxide solar cells this requirement is less stringent, since the thickness of the absorbing layer can be increased up to several nanometers, when covering a nanowire array, getting almost full light absorption in photoanodes with a typical nanowire length around 1e3 mm [39]. However, conformal coverage of an oxide nanowire array is rather hard to be obtained, and typically the p-type layer decorates the nanowire array either as a compact film [23] or as a top layer (Fig. 6.7) [22].

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Figure 6.7 (A) Top view (main frame) and cross-sectional (inset) SEM image of ZnO nanowires. (B) Cross-sectional SEM image of the composite system after the deposition of the Cu2O layer. (C) Current densityevoltage curve of the solar cell under 1 sun irradiation (AM1.5G, 100 mW cm2). Reproduced with permission from T.J. Hsueh, C.L. Hsu, S.J. Changa, P.W. Guo, J.H. Hsieh, I.C. Chen, Cu2O/n-ZnO nanowire solar cells on ZnO : Ga/ glass templates, Scripta Mater., 57 (2007) 53e56. Copyright 2007 Elsevier.

In order to prevent the formation of a compact and thick film, which can hardly transport charges in the operating device, we recently proved the possibility to obtain coreeshell nanowire coaxial heterostructures, in which a thin ptype layer conformal covers a ZnO nanowire [40]. Chemical vapor deposition can be used to obtain a uniform coverage of a p-type layer, with the possibility to tune its thickness and the consequent light absorption of the electrode (Fig. 6.8E). The presence of nanowire geometry induces strong scattering effect, which results in highly reduced transmittance of the ZnO nanowires, compared to ZnO thin film. It is also clear from the electrical measurements that an operating pen junction is established at the ZnO/Cu2O interface, which fastens charge collection under light irradiation. Unfortunately, the obtained open circuit photovoltage (Voc) under 1 sun irradiation is rather low (below 200 mV), so the pen junction can operate as an efficient self-biased photodetector, but its efficiency as a solar cell is far below 1%. The issue of the optimization of nanowire solar cells in terms of light absorption and charge transport is shared also with nonoxide nanowire solar cells. A very recent review highlighted different ways offered by bottomup or top-down preparation methodologies to enhance light absorption through photon management and maximize charge collection, exploiting various geometries with lowered material consumption: axial and radial junctions, substrate interfacial and planar junctions, etc. [41]. One important point is the possibility to move from an optimized, yet single nanowire, to an array of highly efficient solar cells. Aiming at improving the PCE in nanowire solar cells, metal oxides can be fruitfully applied, exploiting their good electronic conduction and their

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Figure 6.8 Formation of coaxial coreeshell Cu2O/ZnO nanowires. ZnO nanowires before (A and C) and after (B and D) the growth of a Cu2O thin film. Different layer thicknesses, from 70 to 420 nm, were tested. (E) Transmittance of ZnO NWs with and without the Cu2O layer. Comparison with a ZnO and a Cu2O thin film is also reported. (F) Current densityevoltage curves in dark (solid lines) and under light irradiation (1 sun, AM1.5 G). (GeH) Self-biased photodetection measurements of bare ZnO nanowires (G) and the composite system with 240-nm-thick Cu2O. Reproduced with permission from P. Ghamgosar, F. Rigoni, S.J. You, I. Dobryden, M.G. Kohan, A.L. Pellegrino, I. Concina, N. Almqvist, G. Malandrino, A. Vomiero, ZnO-Cu2O core-shell nanowires as stable and fast response photodetectors, Nanomater. Energy, 51 (2018) 308e316. Copyright 2018 Elsevier.

energy band gap larger than Si in most of the cases. In a recent report, Y. Zhang and coworkers demonstrated the possibility of integrating n-type Si nanowires with a V2O5 and a TiO2 thin layer, respectively, p- and ndoped [42]. The two layers act as the hole- and electron-selective films, respectively. An interesting feature is that the large difference in work function between Si nanowires, and V2O5 enables the onset of a high built-in potential (equal to 0.75 V), suggesting the possibility to obtain large Voc for this kind of nanowire solar cells under sunlight irradiation. Similar strategy for the n-type layer was also applied in a planar Si cell based on hydrogenated amorphous Si, with the requirement to deposit a very thin TiOx layer, about 1 nm thick [43].

6.4 Back contact and alternative structures The nature of the back contact also plays a relevant role in all-oxide PV devices, as demonstrated by Majhi and coworkers [44]. They

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investigated the influence of four back contact materials (Ag, Cu, Ni, and Co) on Co3O4jTiO2 devices by using a combinatorial approach. Numerical simulations were carried out, devoted to investigate the interfacial recombination at the adsorbereback contact interface. This interface is critical for device performance, and the exciton recombination can be modulated by band bending, surface charge transfer rate, and density of surface states. From their investigation, the metallic back contact mainly affects the density of surface states, with relevant difference among the investigated materials (the highest density of surface states was found in case of Co contact, while Ag featured the lowest one). The correlations found between the nature of the back contact, the thickness of TiO2 layers used as electron transport material, and the cell working parameters are shown in Fig. 6.9. To deal with the recombination at the back contacteadsorber the application of a thin MoO3 layer was also taken into consideration by fabricating devices featuring the following scheme: TiO2jCo3O4jMoO3jAu [45]. The authors observed a systematic enhancement of both VOC and JSC, as shown in Fig. 6.10.

Figure 6.9 Plot of the Voc (top row), Jsc (center row) as a function of absorber layer thickness for four different back contacts. Each plot represents cells with a different back contact material: (A) Ag, (B) Cu, (C) Ni, and (D) Co. The bottom row shows the dependence of the fill factor (FF) on TiO2 (electron conducting layer) thickness; the FF increases with TiO2 thickness. Absorber layer thickness and the back contact material do not show significant effects on FF. Reprinted from K. Majhi, L. Bertoluzzi, D.A. Keller, H.N. Barad, A. Ginsburg, A.Y. Anderson, R. Vidal et al., Co3O4 based all-oxide PV: a numerical simulation analyzed combinatorial material science study, J. Phys. Chem. C, 120 (2016) 9053e9060. with ACS permission. Copyright ACS 2016.

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Figure 6.10 Plots of the (a) Voc, (b) Jsc, (c) FF, (d) IQE, and (e) Pmax as a function of absorber layer thickness for libraries with (green solid dots) and without (blue solid dots) the MoO3 hole-selective contact. Reprinted from K. Majhi, L. Bertoluzzi, K.J. Rietwyk, A. Ginsburg, et al., Combinatorial investigation and modelling of MoO3 hole-selective contact in TiO2jCo3O4jMoO3 all-oxide solar cells, Advanced Materials Interfaces, 3 (2016) 1500405. with Wiley permission. Copyright Wiley 2016.

Concept behind the application of MoO3 thin layer (50 nm) is to support the hole extraction, which is known for being one of the major limiting factors in all-oxide PV. Authors observed that MoO3 is capable of acting as both “hole-selective” contact and passivating layer for the Co3O4 adsorber layer, thus reducing the recombination at the interface with the back contact (Au in this case). Ferroelectric materials have also been investigated as PV materials. A diodelike and photovoltaic effect was indeed observed in BiFeO3 crystals [46]. A zero-bias photocurrent density was recorded under green (532 nm) and red (650 nm) light illumination. The current intensity was higher in the case of green illumination (around 7 mA/cm2) than under red light (in the range of nA/cm2), indicating that photogenerated carriers across the bulk optical gap (around 2.5 eV) dominate the PV effect. The mechanism behind the observed PV effect was elucidated by Seidel and coworkers [47]: when the illumination is strong enough, domain walls between regions of different electrical polarization function as a nanoscale generator of PV current. Charge separation is efficient at the walls, so that electrons and holes are accumulating at different sides, thus reducing the recombination rate at a local level.

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6.5 Conclusions In summary, all-oxide solar cells represent, in principle, a viable alternative for the development of low-cost, nonvacuum-based platform for solar cell production. Unfortunately, at present, the overall performances keep below the expectations. Intrinsic fast and significant recombination of photogenerated carriers possibly represent the main limit of these devices: as discussed, this process occurs in several parts of the cells, thus determining electrical losses that heavily affect the efficiency of solar energy conversion. At present, Cu2O is the most promising absorber material, but more efforts should be put in looking for alternative p-type semiconducting metal oxides. On the other hand, the electron transport materials, in most cases TiO2 or ZnO, also present heavy charge recombination, which is known as the main factor limiting their capability to properly transport the required carriers through the circuit. Investigations have also pointed out the critical role played by the back contact: it is still unclear what the best material may be, in order to reduce recombination at the interface adsorbereback contact. Literature overview eventually highlights the potential hold by these devices, but also points out to the serious problems encountered by the researchers in fabricating cells featuring reasonable performance. On the other hand, it seems here worth noting that the scientific community active in this field is rather small, possibly due to the complexity of materials preparation and cell assembly. Indeed, in spite of the possibility to apply vacuumfree and cheap synthetic methods for the synthesis of the required materials, and of the apparently easiness of device assembly, all the studies have remarked that both aspects are less trivial than expected. This suggests that more effort is needed if we want to access materials with improved characteristics and, in the future, devices with decent performance.

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CHAPTER SEVEN

Simulations of conventional and augmented types of solar cells Alexander Quandt1, Itumeleng Mokgosi1, Robert Warmbier2 1

School of Physics, Centre of Excellence in Strong Materials and Materials for Energy Research Group, University of the Witwatersrand, Johannesburg, South Africa Department of Physics, University of Johannesburg, Auckland Park, South Africa

2

7.1 Introduction In order to fabricate highly efficient photovoltaic devices, one can fall back on a whole variety of basic materials, ranging from classic inorganic materials like crystalline Si or GaAs, over perovskites, CdTe and Cu(In,Ga)(Se,S)2 (CIGS), to thin film systems made of organic materials or inorganic materials like amorphous Si:H. The advantages and disadvantages of these materials are well described in a recent review article [1]. Beyond the choice of the basic materials, there is also the problem of designing and fabricating an efficient solar cell. State-of-the-art design of highly efficient solar cells seems to point toward increasingly complex systems, as outlined in Ref. [2]. Among the various strategies to increase the photon flux into a solar cell, there are specially designed antireflective coatings, photonic crystal layers, and plasmonic nanoparticles [3,4]. Furthermore, photons can be harnessed in the ultraviolet (UV) and infrared (IR) spectrum using up- and downconversion [5]. These photons, otherwise, would either not be absorbed (IR) or their absorption (UV) would lead to considerable losses due to the thermalization of excited electrons toward the bottom of the conduction band. We will discuss these frequency conversion approaches in more detail below. With such a large variety of basic materials and enhancement strategies, it makes sense to shift some of the more complex design problems from the lab onto a computer and thus perform solar cell device simulations [6]. This will be the topic of the present chapter. We will start this chapter with an elementary discussion of pen junctions, where we describe the simplest modeling approach based on the diode equation [7,8]. This basic approach already contains important features of Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00007-0

© 2020 Elsevier Ltd. All rights reserved.

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more advanced simulation techniques. Therefore, we can use the diode equation to discuss general strategies, which will increase the efficiency of standard inorganic solar cells. In a subsequent section, we will give a short introduction into state-ofthe-art solar cell device simulations, where we essentially follow [6]. Our focus will be on a basic set of equations that describe the physics of a solar cell. We will also discuss the numerical methods to solve these equations, where we will have a special attention on the materials parameters that enter in these simulations. It turns out that the materials parameters for solar cell device simulations can all be obtained from ab initio simulations based on density functional theory and related methods. The use of ab initio methods opens the exciting possibility to develop new types of solar cells entirely in silico, and virtually by starting from the atomic structures of the basic materials. Such an approach is not part of the standard literature on solar cell device simulations, and therefore it might be of particular interest for some of our readers. In the final section, we will discuss the implementation of frequency conversion layers into state-of-the-art solar cell device simulations. These device simulations are supposed to be carried out alongside experimental studies to improve the efficiency of commercial types of solar cells, where the frequency conversion is implemented by the deposition of glass ceramic layers containing rare earth ions. The implementation of frequency conversion processes is also a very recent development in solar cell device simulations, and we will present preliminary results from some of the first studies that have tackled this challenging problem. Note that each of these novel and unconventional approaches to solar cell device simulations is still work in progress, and it is fair to say that we might have only seen the beginning of entirely new developments in this field up to now. Therefore, we will close our chapter not only with a short summary of our major findings but also with an outlook on possible new directions and future opportunities in the field of solar cell device simulations.

7.2 About pen junctions, diodes, and solar cells 7.2.1 Basics of pen junctions Inorganic solar cells are based on semiconductor pen junctions. The corresponding basic layout of an inorganic solar cell is shown in Fig. 7.1, together with the electrodes that are used to extract current, and thus power,

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Figure 7.1 Layout of a typical inorganic solar cell.

from such a device under solar illumination. Another important feature of modern solar cells are antireflective coatings, which scatter light preferentially into the active region of a solar cell, and thus increase the absorption of photons and the generation of a photocurrent. A typical pen junction consists of two doped semiconductors, as shown in Fig. 7.2. The semiconductor on the p-side is doped with acceptor atoms, which have one valence electron less than the atoms of the bulk material. This leads to an excess of very mobile holes. Such materials are commercially fabricated, e.g., by doping crystalline silicon with boron atoms. On the nside the bulk semiconductor is doped with donor atoms, which have one valence electron more than the atoms of the bulk material. This leads to an excess of very mobile electrons. A typical example of such an n-layer would be bulk silicon doped with phosphorous atoms. At the interface between the two materials, electrons from the n-side (majority carriers) diffuse into the p-side (minority carriers), where they annihilate with the holes. Likewise, holes diffuse from the p-side into the n-side, where they annihilate with the electrons. Therefore, close to the interface we will see the formation of a whole region devoid of any free charge carriers, and where the net charges of the fixed dopants start to repel free charge carriers from the other side of the junction. This system eventually stabilizes itself by forming a depletion layer as indicated by the red arrows in Fig. 7.2. A typical pen junction has rectification properties, which are sketched in Fig. 7.2B and C. Therefore, it may serve as a diode. To understand this particular feature of a pen junction, it is important to know that only the

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Figure 7.2 pen junction. The semiconductor on the p side is heavily doped with acceptor atoms (indicated by large gray minus sign), which leads to an excess of very mobile holes (indicated by small white plus signs). The semiconductor on the n side is heavily doped with donor atoms, which leads to an excess of very mobile electrons. The red arrows indicate the size of the depletion layer, which is formed at the interface between both semiconductors. (A) Junction without applied voltage (zero bias). (B) Forward bias. Note the shrinking of the depletion layer. (C) Reverse bias. Note the expansion of the depletion layer.

majority carriers (i.e., electrons on the n-side and holes on the p-side) have high enough mobilities to flow through these parts of the pen junction, whereas the minority carriers (i.e., electrons on the p-side and holes on the n-side) can only migrate over very short distances. In forward bias (Fig. 7.2B), there is a positive terminal connected to the p-side, and a negative terminal connected to the n-side. Under these conditions, additional majority carriers are pushed into the depletion layer from both sides, which decreases the size of the layer, and thus allows for a constant influx of minority carriers. Although the latter cannot migrate very far, their influx is compensated by a macroscopic flow of the majority carriers, which leads to a net current through the device. In reverse bias (Fig. 7.2C), a negative terminal is connected to the p-side and a positive terminal is connected to the n-side. In such a situation the majority carriers are pulled away from the interface, which leads to an expansion of the depletion layer, and to the formation of such a huge repulsive barrier at the interface, that the influx of minority carriers basically stops.

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Only a very small current may still flow under such conditions, which is called the saturation current Js (here and in the following we work with current densities J rather than currents).

7.2.2 A diode model for solar cells Before we describe the behavior of a pen junction under applied bias in a quantitative fashion with the help of Shockley’s diode equation [9], we want to briefly discuss the working of a typical inorganic solar cell. When a pen junction in forward bias is exposed to light, the incoming photons will be absorbed by the bulk semiconductor due to the creation of bound electrone hole pairs called excitons, which quickly dissociate due to the strong electrostatic fields in the depletion layer. The resulting free charge carriers preferentially move into those regions, where they have the highest mobilities (majority carriers). This movement leads to a net current flow in a direction opposite of the current flow in a diode under forward bias, and the system is actually generating power. The resulting current is called the solar current Jl . By effectively separating the photogenerated free charge carriers, the pen junction basically acts like a membrane [10]. We have already identified the diffusion of minority charge carriers as the driving force behind the current in a diode. In the case of electrons drifting into the p-side of the junction, we observe the following current density Jn : Jn ¼ e$v$ðn np Þ

(7.1)

Here e is the elementary charge, and v is the group velocity of the electrons, which is related to the band structure EðkÞ of the semiconducting bulk material via v ¼ dEðkÞ dk . Furthermore np is the concentration of electrons (minority carriers) far away from the depletion layer on the p-side of the junction, and n is the local concentration of electrons along the diffusion path. We have also discussed before that a voltage V can be used to regulate the diffusion of minority carriers across the junction, and the temperature T will influence this diffusion process as well. Therefore, it makes sense to assume the following Boltzmann law for n: eV

n ¼ np ekB T

(7.2)

where kB is Boltzmann’s constant. Based on this assumption, we arrive at the following expression for the electronic current density:
eV  (7.3) Jn ¼ e$v$np $ ekB T  1

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A similar reasoning applies to the current density caused by the diffusion of holes into the n-side of the junction, which is:
eV  (7.4) Jp ¼ e$v$pn $ ekB T  1 where pn is the concentration of holes (minority carriers) far away from the depletion layer on the n-side of the junction. The total current density J is simply the sum of the currents due to the diffusion of electrons and holes:
eV  
eV J ¼ e$v$ðnp þ pn Þ$ ekB T  1 hJs ekB T  1 (7.5) which defines the saturation current density Js . In the case of a solar cell the pen junction is exposed to light, and therefore we have to add a solar current density Jl to the diode equation, which leads to 
eV J ¼ Js ekB T  1  Jl (7.6) Note the change in sign compared to the saturation current density Js , which was discussed above. The solar current density Jl can easily be 10 orders of magnitude larger (or even more!) than the saturation current density Js , which is important to know when it comes to the design of a working photovoltaic device. A simple model for the solar current density Jl in crystalline silicon would be the following: Jl ¼ e$g$nSi $d (7.7) Here g is the generation rate of electronehole pairs per Si atom, nSi is the atomic density of bulk Si, and d is the size of the depletion layer. Assuming a typical generation rate of g ¼ 0:01 1s per Si atom, an atomic density of nSi ¼ 5$1028 m13 and a depletion layer of size d ¼ 104 m, we obtain a solar current density of Jl ¼ 8$103 mA2 . A typical solar current density in a conventional Si solar cell is about 500 mA2 , which is an order of magnitude smaller. The reason for the failure of such a simple model is that it neglected recombination effects, which effectively reduce the generation rate g. These effects are included in a full-fledged device simulation and will be discussed in a subsequent section. Note that the product G ¼ g$nsi is essentially the same as the frequencyand position-dependent generation rate Gðu; xÞ used in solar cell simulations, which can be determined from Beer’s law [6]: Gðu; xÞ ¼ aðuÞ$fðuÞ$eaðuÞ,x

(7.8)

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Here aðuÞ is the absorption coefficient of the bulk semiconductor, which can be measured experimentally or determined using ab initio methods, and fðuÞ is the incoming photon flux. For a typical solar irradiaW and a typical photon energy of E z1019 J, we have an tion of Isol z1000 p m2 1022 . With an absorption coefficient of estimated photon flux of about fzs,m 2 1027 , which is consistent az105 m1 , we obtain a generation rate of about Gzs$m 2 with our previous assumptions only at the surface of the solar cell, but not within the bulk, where it actually drops according to Eq. (7.8). Therefore, we may conclude that our chosen generation rate of g ¼ 0:01 1s is somewhat on the optimistic side, and that this number most likely refers to much stronger illuminations by a laser rather than by the sun itself. In Fig. 7.3, we show the typical current density J versus voltage V relations for a diode and a solar cell, which follow from Eqs. (7.5) and (7.6). Both systems are nonlinear devices. Note that the right-hand side of the figure corresponds to positive voltages V , where both devices are in forward bias, and the left-hand side of the figure corresponds to negative voltage, which is the situation for reverse bias. We notice that the solar current density pulls the characteristic curve of a diode in the negative direction of the current density because the solar current density Jl is much larger than the saturation current density Js . The so-called open circuit voltage of the solar cell Voc can be obtained from Eq. (7.6) by setting J ¼ 0, which leads to     kB T Jl kB T Jl Voc ¼ (7.9) þ1 z ln ln e e Js Js

Figure 7.3 Current density (J) versus voltage (V ) characteristics of a diode (blue) and of a solar cell (red). Voc is the open circuit voltage of the solar cell, Pm is the power per area at the maximum working point (black dot), Jl is the light current density and Js is the saturation current density.

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With a typical ratio of JJsl z1012 and kBeT ¼ 0:0258 V , we obtain open circuit voltages of about Voc z 0.7 V, which is what we usually measure for crystalline Si under solar illumination. The open circuit voltage Voc is often tabulated to characterize a given type of solar cell device, together with the short-circuit current density Jsc ¼ JðV ¼ 0Þ. For a solar cell in forward bias, but below the open circuit voltage V , the power per unit area P ¼ J$V is negative, which means that the solar cell actually produces power. The maximum power Pm per unit area can be obtained at the so-called maximum working point ½Vm ; Jm , where Pm ¼ Jm $ Vm . The corresponding voltage Vm can be determined from which amounts to

dPðVm Þ dV

¼ 0,

Jl   1þ ! kB T kB T Jl Js Vm ¼ ¼ Voc  (7.10) ln ln þ 1 eVm e e Js 1þ kB T This equation obviously involves a nonlinear relation to determine Vm . As Jl [Js , we find that Vm < Voc , as shown in Fig. 7.3. The corresponding current density Jm can be obtained by inserting Vm into Eq. (7.6). The power per unit area Pm corresponds to the yellow rectangle shown in this figure.

7.2.3 Detailed analysis of the diode model We now use Eq. (7.6) to design solar cells and to predict their properties. This is shown in Fig. 7.4. In the case of Fig. 7.4A, we have changed the solar current density Jl from 0 to 500 mA2 in steps of 25 mA2 , while keeping the saturation current density at Js ¼ 5,108 mA2 . According to Eq. (7.9), this leads to an open circuit voltage of Voc ¼ 0:59 V for a solar current density of Jl ¼ 500 mA2 , which is what we observe for the lowest curve with the biggest power output in Fig. 7.4A. It is obvious that the efficiency increases in a nonlinear fashion with increasing solar current density, as can be seen from the changes in location of the maximum working point. For the curve with the maximum power output we have a maximum power per unit area of Pm ¼ 233:17 W m2 . We can estimate the efficiency h of such a solar cell by using the following equation: power out jPm j $100 % ¼ $100 % (7.11) Isol power in With an estimated solar irradiation of Isol ¼ 1000 W m2 , we obtain an efficiency of about h ¼ 23:3 %, which is pretty good. However, there are h¼

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Figure 7.4 Current density versus voltage characteristics (A) for a series of solar cells with different light current densities. Note the shift of the maximum working point with increasing light current. (B) A series of solar cells with different saturation current densities. Note that the increase in power is marginal, even after the saturation current density has been scaled by an unrealistic factor of 107 as compared to (A).

many processes that may reduce the efficiency of a solar cell, starting from an inefficient injection or extraction of charge carriers at the electrodes [6], which is completely absent from our model. Therefore, the calculated efficiencies from such a simple model must always be taken with a grain of salt. However, we clearly see from Fig. 7.4A that any increase in the solar current density should always be a promising strategy to increase the power output of a solar cell, and according to Eqs. (7.7) and (7.8) we can achieve such an increase in various ways. The first strategy would be to increase the photon flux fðuÞ, e.g., by using solar concentrators, suitable antireflection coatings, plasmonic nanoparticles, or patterned electrodes [3]. Furthermore, we can use frequency conversion techniques to convert light from the UV and IR, with the goal to increase the photon flux in a frequency range where a conventional solar cell is most efficient. This will be discussed in more detail toward the end of this chapter. Alternatively, we may try to increase aðuÞ and thus the absorption of incoming photons, but not to the point that no light will be able to reach the depletion layer any more. In fact, it would be advantageous to increase the absorption only very close to the depletion layer, e.g., by using plasmonic nanoparticles. A third strategy would be to increase the size of the depletion layer, and thus the probability for a photon to create free charge carriers after dissociation. A forth strategy would be to reduce the recombination rates, and thus increase the net generation rates. All of these strategies can be examined based on Eq. (7.6).

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Now suppose that someone comes up with the idea to increase the saturation current density Js at fixed solar current density Jl ¼ 500 mA2 . According to Eq. (7.5), this could be achieved by modifying the band structure and thus the group velocity v, or by increasing the concentrations of minority carriers far away from the depletion layer. We have simulated such an approach in Fig. 7.4B, where we have increased the initial saturation current density from Js ¼ 5$108 mA2 in huge steps of 0:5 mA2 . Note that these huge steps were necessary in order to see any noticeable effect. It is quite obvious from Fig. 7.4B that an increase in the saturation current density has only a very small effect on the power produced by the corresponding solar cell. We can check the consistency of Fig. 7.4B by looking at a curve with a saturation current of about Js ¼ 0:5 mA2 . According to Eq. (7.9), this should lead to an open circuit voltage of about Voc ¼ 0:178 V , which is what we see in Fig. 7.4B. Being the scenario with the largest power output, we find that Pm ¼ 56:97 W m2 , which changes only very little for other saturation current densities. The corresponding efficiency according to Eq. (7.11) is h ¼ 5:69 %, which is not very good. By tampering with the saturation current density, we have obviously chosen a very bad strategy to increase the efficiency of a solar cell. In the following sections, we will describe more advanced methods to simulate solar cells [6], which properly take into account the scattering of light by different layers inside the solar cell, provide models for the extraction and injection of charge carriers at the electrode, and take into account recombination and nonradiative processes inside a solar cell. Furthermore, these advanced simulation methods also comprise proper models for the solar photon flux over a large frequency range.

7.3 Solar cell device simulations Now we move on to more advanced simulations of solar cell devices. From a numerical point of view, the primary challenge consists of solving a system of three coupled nonlinear differential equations [6]. These basic equations are the Poisson equation, the continuity equation for the holes, and the continuity equation for the electrons. In the framework of solar cell simulations, one usually solves these equations under steady-state conditions. In order to obtain realistic results, it is also necessary to impose physically meaningful boundary conditions. Thus, depending on the details of the physical processes involved, and depending on the complexity in the

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layout of a given photovoltaic device, numerical solar cell device simulations can be quite sophisticated. There are several advanced device simulation packages such as AMPS1D [11,12], SCAPS [13], and GPVDM [14], which essentially solve the basic set of coupled differential equations under appropriate boundary conditions and produce quite reliable IeV characteristics. The various numerical techniques used to solve the basic equations are detailed in the respective manuals of these simulation packages. AMPS-1D [11], for example, uses a discretization approach, where the NewtoneRaphson method is used to solve a set of 3(Nþ1) algebraic equations given in a suitable matrix format. This set of governing algebraic equations basically results from dividing the solar cell device structure into N slabs. Further technical details are very well described in textbooks like Solar Cell Device Physics [6], and in the following we will only present the general theoretical framework of state-of-the-art solar cell simulations. Let us start with a description of the basic set of coupled differential equations. The role of the Poisson equation is to associate the carrier concentrations with the electrostatic potential that is present throughout the whole system. The continuity equations, on the other hand, describe the change of the carrier densities under different applied biases. These equations will essentially keep track of all the electrons from the conduction band and of all the holes from the valence band of the bulk semiconductor materials. They also take into account the presence of additional free charge carriers introduced by the doping of the bulk materials. The Poisson equation and the continuity equations are solved in one characteristic dimension, which we will denote by x. The incoming light is supposed to impact the solar cell at right angles to the top layer, and then it bounces back and forth along this direction until it hits the bottom contact of a typical solar cell device. The resulting set of equations in the steady state is [15]: dE e ¼ ðp  n þ ND  NA Þ dx ε

(7.12)

d d2 n ðn$EÞ þ Dn 2 ¼ eðR  GÞ dx dx

(7.13)

d d2 p ðp$EÞ  Dp 2 ¼ eðR  GÞ: dx dx

(7.14)

mn mp

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Here, ε is the permittivity of the bulk semiconductor, E is the electrostatic field across the system, ND is the concentration of donor atoms on the n-side, and NA is the concentration of acceptor atoms on the p-side. R and G denote the net recombination rate and the carrier generation rate within the semiconductor material. p and n are the free carrier concentrations for the holes and the electrons, respectively. Note that we have shown some of the symbols in bold face, which should indicate that all of these entities are based on materials parameters that could in principle be determined using ab initio methods. We will do the same for other equations discussed in this section, before we discuss these fundamental ab initio materials parameters in more detail as part of the next section. It turns out that the carrier concentrations depend on the corresponding Fermi levels. Assuming parabolic model bands for electrons within the conduction band and the valence band, and based on the corresponding densities of states NC and NV at temperature T , one may derive the following general expressions for the carrier concentrations [16]:   EF  E C n ¼ NC F1=2 exp (7.15) kT   E V  EF p ¼ NV F1=2 exp ; (7.16) kT where EF is the Fermi level of the bulk semiconductor material, EC marks the bottom of the conduction band, and EV marks the top of the valence band. The difference between EC and EV is the electronic band gap E g : F1=2 in Eqs. (7.15) and (7.16) denotes the FermieDirac integral of order ½ which can either be computed numerically or taken from tabulated data. Note that as soon as light starts to enter the solar cell, the material is driven out of equilibrium, and the Fermi level EF of the bulk material splits into two separate quasi-Fermi levels EFn;p for either the electrons or the holes (socalled quasi-equilibrium). This means that the electrons themselves are still in equilibrium with each other, and so are the holes, but the two systems of free charge carriers are no longer in equilibrium with each other. The quasiFermi levels EFn;p will depend on the position x, and therefore the left-hand dE side of Eq. (7.13) can also be written as emn dxFn . The same applies to the lefthand side of Eq. (7.14).

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Typical expressions for effective densities of states NC and NV in Eqs. (7.15) and (7.16) are given by Ref. [16]:  3=2 2p$m*n $kB T NC ¼ 2 (7.17) h2  3=2 2p$m*P $kB T NV ¼ 2 : (7.18) h2 From these simple models, we see that the effective densities of states depend on the effective masses m*n and m*p of the charge carriers, and thus on the details of the electronic band structure of the bulk material. The carrier mobilities mp and mn also play an important role in Eqs. (7.12)e(7.14). They are related to the diffusion coefficients Dn and Dp through the Einstein relationship Dm ¼ kBeT [16]. The generation rate G in Eqs. (7.13) and (7.14) is given by Ref. [17]: Z GðxÞ ¼ ð1  Sf Þ ð1  rðlÞÞfðlÞaðlÞeaðlÞ,x dl : (7.19) l

Here Sf is a shadowing factor, which is related to the wire grid that connects individual solar cells in a panel. The term rðlÞ denotes the reflectance from the surface, and aðlÞ is the absorption coefficient as a function of wavelength using the relationship Zu ¼ hc=l. The term fðlÞ denotes the incident photon flux, where it is assumed that sunlight enters the solar cell perpendicular to the top surface at x ¼ 0. Eq. (7.19) is basically a modified and integrated version of the frequency-dependent generation rate discussed in Eq. (7.8), which describes the rate at which incoming photons of a given wavelength l will generate pairs of free charge carriers. For the simulations, the incident photon flux FðlÞ refers to a variety of models for the solar spectrum at the surface of the Earth, and the absorption coefficient aðlÞ describes the basic optical properties of the bulk semiconductor material used for the solar cell. It is shown in bold face, which means that it can be obtained from ab initio methods, as discussed in the next section. The net recombination rates R in Eqs. (7.13) and (7.14) consist of three basic recombination mechanisms [6]: recombination through defects (traps) within the band gap, Auger recombination, and band-to-band (radiative) recombination. Recombination through traps is a two-step process, where electrons relax from the conduction band to an intermediate state

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introduced by a defect. Afterward they further relax to the valence band and annihilate with a hole, which is assisted by phonons. Auger recombination is a process where the transition energy is passed onto another charge carrier in either the valence or conduction band, which relaxes thermally afterward. Radiative recombination is the optical generation process in reverse and is more prevalent in semiconductors with direct band gap. In this process, the energy of the decaying electron is transferred to an emitted photon [17]. For all of these recombination processes, we can define associated carrier lifetimes se;h for the electrons and the holes, which are related to the net recombination rate R by Ref. [15]. se ¼

Dn R

(7.20)

Dp : (7.21) R The two terms Dn and Dp describe disturbances of the respective carrier concentrations from their equilibrium values. Due to the sudden termination of the crystal lattice at the contacts, we find a huge concentration of defects within that particular region, caused by the surface states on both sides of the interface. Detailed models for the resulting surface recombination rates RS can be found in Refs. [3,17]. Let us summarize some of our main findings: we have seen in this section that the design of basic semiconductor solar cell boils down to a layered cake model, where key physical entities only vary along one characteristic direction x. The main features that will determine the size of the generated photocurrents are certainly the two doped bulk semiconductors. Other important features are the two contacts, where the free charge carriers are extracted, and an antireflection layer on top of the solar cell, which will help to reduce reflection losses. The modeling of the corresponding device performance depends on a relatively small number of parameters [18]: 1. The dopant concentrations ND and NA on the n- and p-side of the junction. Their concentrations are basically controlled during synthesis of a real device, and we must use the same values for the corresponding device simulations. 2. The carrier mobilities mn and mp and the related diffusion coefficients Dn , Dp , which characterize carrier transport within the device. sp ¼

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3. Materials properties related to the electronic structure EðkÞ of the bulk semiconductor, like the band gap energy Eg , the top of the valence band EV , and the bottom of the conduction band EC . The complex electronic band structure of a bulk semiconductor can often be approximated by a parabolic band structure, and these parabolic models are typically characterized by parameters like the effective masses m*n and m*p . 4. Optical materials properties like the frequency (or wavelength)dependent permittivity ε, as well as the related absorption coefficient a. It must be pointed out that the path of a single photon inside a solar cell can be quite erratic, which means that basically all of the optical materials parameters are needed to describe the fate of such a photon. Here the absorption coefficient a of the bulk semiconductor plays a particular role, because it determines the generation rate of free charge carriers. 5. In order to model the main recombination processes, we need to know other materials parameters like the local concentrations of traps, and we need to know order of magnitude values for the typical carrier lifetimes se and sp in a bulk semiconductor. Alternatively, one could use simple atomistic models to specify recombination cross sections, and then apply Beer’s law.

7.4 Ab initio materials properties At the end of the last section, we gave a list of key simulation parameters for a typical solar cell device simulation. In the following, we will show that all of the materials parameters described in items 2e4 can in principle be taken from first principles simulations of the corresponding bulk materials, based on density functional theory (DFT). This means that we can simulate a solar cell device without ever touching a real material or even entering a laboratory. Let us start with item 3, which refers to materials properties related to the electronic structure of bulk semiconductors. Although it is known that the electronic band gap Eg is generally underestimated by DFT methods, the local curvature of the bands is usually quite reliable, which is important for calculation of the effective masses m*n and m*p . These are indeed calculated from the curvature of the conduction and valence band minimum and maximum, respectively. More specifically, the effective masses of a band n at point k0 in directions i, j are given by Ref. [19]:

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m*ij

" #1 Z2 v2 En ¼ 2 vki vkj

(7.22)

k¼k0

Next, we want to discuss item 2, which refers to the carrier mobilities mn and mp and the related diffusion constants. In general, the carrier mobility e m ¼ * srelax (7.23) m is directly proportional to the relaxation time srelax , which is the average time between collisions of the charge carriers with the lattice in a simple Drude model. The relaxation time is also identical to the imaginary part of the electron self-energy Snk at the center of the Brillouin zone. Therefore, a typical relaxation time may be obtained from an ab initio solution of the electron self-energy equation, which is associated with the electrone phonon interaction by Refs. [20,21]: " X   nqn þ fεmkþq 2 v Snk ¼ 2 wq gmn ðk; qÞ εnk  εmk þ q þ uqn  ih qv # nqn þ 1  fεmkþq . (7.24) þ εnk  εmk þ q þ uqn þ ih n ðk; qÞ are the electronephonon matrix elements given by: Note that gmn     v ðk; qÞ ¼ jmkþq vqn V jnkþq . (7.25) gmn

The electronic wave functions jnk of the nth band with wave vector k, and with corresponding band energies εnk and Fermi occupations fnk , can be taken directly from DFT ground-state calculations. The frequency of the vth phonon branch at wave vector q is denoted by uqv , with nqv being its Bosee Einstein occupation number. Furthermore, vqv V is the derivative of the selfconsistent potential associated with the vth phonon at q. The numerical parameter wq in Eq. (7.24) is the weight of the point q in the Brillouin zone summation. The electronephonon matrix elements and the phonon wave functions can be determined using density functional perturbation theory [22], which is described in Refs. [23,24]. Let us finally discuss item 4, which refers to optical materials p properties. ffiffiffiffiffiffiffiffiffiffiffiffiffi  The frequency-dependent absorption coefficient aðuÞ ¼ 2u J εM ðuÞ c can be obtained from the macroscopic dielectric constant εM ðuÞ. The latter can be calculated to a certain accuracy using linear response theory within the DFT framework [25]. But such a procedure tends to underestimate

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the absorption at the absorption edge, and it does not consider excitonic effects all together, which means that this approach will neglect the creation of bound electronehole pairs by the incoming photons. The exact response of the material to optical excitations is described by the four-point BetheeSalpeter equation (BSE). To make the BSE solvable for more than just the most trivial model systems, several approximations and simplifications have to be made [26,27], which lead to an effective two-point response function with only real eigenvalues similar to DFT. ðn n Þ With the BSE eigenstates, eigenvalues El and coefficients Al 1 2 for the particleehole pairs that contribute to this mode, the (complex) macroscopic dielectric function is given by: P   iq:r  "  ðn1 n2 Þ 2 #   n2 Al  X  n1 n2 n1 e εM ðuÞ ¼ 1  lim ðnG¼0 Þ  . (7.26) q/0 u  E þ ih l l The labels n refer to whatever one-electron basis is used to expand the BSE states in. Usually, these labels refer to DFT states. Furthermore, vG¼0 ðqÞ is the Coulomb kernel in reciprocal space, and h is a suitably chosen broadening parameter. In Table 7.1, we show a typical set of ab initio parameters calculated with Quantum Espresso [28], which are in good agreement with experimental values. Differences between experimental and numerical data for the effective masses and mobilities can of course stem from limitations of the DFT framework, but also from different treatment of light and heavy holes/electrons between experiment and numeric data. We used these parameters to simulate the performance of the corresponding solar cell, which is shown in Fig. 7.5, together with the plot of Table 7.1 A set of ab initio materials parameters for crystalline Si, which were determined using the Quantum Espresso ab initio program package [29]. The corresponding wavelength dependent absorption coefficient aðlÞ is shown in Figure 8.5a. Materials parameter Experimental value Numerical value

Effective electron mass, m*n Effective hole mass, m*p Electron mobility, mn Hole mobility, mp Band gap, E G

k: 0:92m0 ; t : 0:19m0

k: 0:90m0 ; t : 0:19m0

l : 0:15m0 ; h : 0:53m0 1350 cm2 V 1 s1 480 cm2 V 1 s1 1:12 eV

l : 0:09m0 ; h : 0:63m0 964 cm2 V 1 s1 218 cm2 V 1 s1 1:24 eV

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Figure 7.5 Simulation of a Si-based solar cell using the WxAMPS software package and the ab initio parameters of Table 7.1. (A) Shows a plot of the corresponding ab initio absorption coefficient aðlÞ for bulk Si. (B) Comparison of JeV characteristics based on experimental data (orange) and on ab initio data (blue). We obtain an efficiency of 2.47%, which compares well with the efficiency of 2.69% for a prepacked parameter set provided by the software package. BSE, BetheeSalpeter equation.

a typical ab initio absorption coefficient aðlÞ calculated at BSE level with the Yambo ab initio code [29]. The results are almost identical to the results for a reference parameter set provided by the software package WxAMPS [30], which is based on experimental data. These simulations demonstrate that we have all of the numerical tools at hand to develop novel types of solar cells entirely on a computer. We may even start our simulations from first principles, based on a quantum mechanical description of the atomic structures related to the constituent materials. The numerical costs of ab initio simulations scale usually at least with the square of the number of valence electrons in the bulk material. Quantities related to the DFT band structure demand little resources and can be computed even on standard quad core work station computers within minutes or hours for small unit cells ((10 atoms). Calculations involving electronephonon coupling are highly demanding due to the involved integral over the Brillouin zone. Those can easily take weeks on a single 16 core compute server even for small unit cells. Dielectric functions based on linear response DFT would take less than a day on such a server for small systems. BSE calculations are much more demanding. They easily take several weeks on one such server, and realistic simulations need to be done on massive parallel high-performance compute clusters.

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7.5 Simulations of augmented solar cells In the introduction to this chapter, we have mentioned some of the strategies to increase the efficiency of conventional solar cells, e.g., by using antireflective coatings, photonic crystal layers, and plasmonic nanoparticles [3,4]. Another approach would be to implement frequency-conversion processes to harness photons outside the spectral range of a typical commercial unit cell [5]. In the following, we want to discuss, how some of these augmentation features could be implemented into the numerical framework outlined up to now. The implementation of a thin layer of plasmonic nanoparticles at the top of the solar cell should lead to a preferential scattering of photons into the active layer of the solar cell, which should also reduce reflection losses due to the use of thin film technologies rather than depositing thick antireflective coating. Furthermore, the presence of plasmonic nanoparticles close to the depletion layer should locally increase the generation of free charge carriers, which works against the various recombination mechanisms discussed before. The easiest way to implement the presence of plasmonic nanoparticles as part of an advanced simulation of an augmented solar cell device would be by modifying the absorption coefficient aðuÞ in Eqs. (7.8) and (7.19) using an effective medium approach. We have already explored this approach and obtained very encouraging results, which showed an increase in efficiency of several percentages for a conventional Si-based solar cell [31], which seems to be in agreement with typical experimental results. In the following, however, we will not delve into the plasmonic approach but will focus on the modeling of frequency conversion layers in standard simulations of conventional Si-based solar cells. More details can be found in Refs. [32,33], but it must be pointed out that these types of simulations are new to the field, and therefore they are still a work in progress.

7.5.1 Physical background In the field of photovoltaics, upconversion (UC) is an interesting optical process by which (several) photons with energies less than the electronic band gap of a bulk semiconductor are converted into higher energetic photons, which can be absorbed quite efficiently by this bulk semiconductor. In the case of photovoltaic devices based on crystalline silicon it is estimated

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that about 20% of the solar spectrum with photon energies lower than the band gap energy could in principle be upconverted by an UC layer situated at the rear of the device and thus improve the overall efficiency of such a conventional solar cell quite considerably [5]. The UC process makes use of a luminescent material (phosphor), which has multiple energy levels with favorable level spacing to emit photons within the optimum frequency range of a given solar cell device. This premise makes lanthanide ions such as Er3þ (trivalent erbium), or Er3þ codoped with ytterbium [34], very attractive for use in UC applications [35,36]. The typical experimental procedure is to dope optical materials like glasses and ceramics with one or more of these activator ions like Er3þ, and often also include sensitizer ions such as Yb3þ as codopants [37,38]. The latter will increase the overall absorption cross section through energy transfer (ET) to the activator and thus improve the overall UC efficiency [38]. This scheme is called energy transfer upconversion, and a typical example is shown in Fig. 7.6. Another frequency conversion technique is downconversion (DC). Here some of the higher energetic photons, which normally generate free charge carriers that will lose most of their energy to thermalization, may be converted into (several) lower energetic photons. These lower energetic photons can generate free charge carriers with very small thermal losses [5]. The corresponding DC layers should be deposited on top of the solar cell.

4 F 4 7/2 H 4 11/2 S3/2

4 Cb2

σ24

3

15

4

Cb1 10

4F 5/2

4F 7/2

1

a Yb3+

C W3 22

W21

0.80 0.98

W4 4

σ02 W2 σ20 W1

F9/2 0.66

4I 9/2 4I 11/2

σ13

2

C2a σba σab Wb

5

0

b

Cb0

0.49 0.52 0.55

I13/2 1.53

Wavelength / μ m

Energy / 103 cm–1

20

C11 4I 15/2

0

0

Er3+

Figure 7.6 Term scheme of a typical Er/Yb upconversion system, which was taken from Ref. [42]. Here the population densities N3 and N4 turn out to be in the right frequency range for the augmentation of standard Si based solar cells.

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The theoretical description of a complex coupling scheme like in Fig. 7.6 is based on rate equations for the population densities Ni of the various energy levels involved. The system is pumped by a flux density fp , which could come from a laser or from sunlight (in the case of a solar cell). As the system is pumped from the ground state of the rare earth ions, the final population densities will depend on the concentration of the rare earth ions in the sample, and therefore these concentrations have to be specified as well. A suitable frequency conversion model has to take into consideration all of the relevant processes involved in the frequency conversion process, such as ground state absorption, spontaneous emission, excited state absorption, stimulated emission, multiphonon relaxation (MPR), and ET. For a very general approach to solve the resulting systems of rate equations see Ref. [37]. The various cross sections, decay rates, and ET rates indicated in a scheme like Fig. 7.6 are often determined from experimental data. Alternatively, some of these coupling constants can be calculated from the JuddeOfelt theory (JO) [39,40]. The most attractive features of JO theory are a sound theoretical background to determine oscillator strengths for absorption and luminescence processes, luminescence branching ratios, excited-state radiative lifetimes, intensity parameters, and transition probabilities [38].

7.5.2 Numerical implementation The simplest approach to include UC and DC into solar cell simulations is to first determine all population densities Nb for states with label b, which will generate photons in the right frequency range of the solar cell. To this end, one has to solve the complex system of rate equations related to term schemes as shown in Fig. 7.6, which is usually done in the steady state. The level b itself decays according to. dNb 1 ¼ Nb sb dt

(7.27)

where sb is its radiative lifetime. It will therefore emit photons at a rate s1b Nb . The corresponding additional photon flux fex ðlÞ for a conversion (glass) layer of size d can be estimated by Ref. [33]: fex ðlÞ ¼

1 Nb d sb

(7.28)

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Once the additional photon flux is known, we can modify the spectral flux models of solar cell simulation packages and otherwise run the simulations as usual, eventually taking into account the additional optical properties of the conversion layer. Such a procedure is easy to implement, and it leads to a proper description of frequency conversion processes in the framework of solar cell simulations. We can also choose a simpler approach by directly modeling the resulting JeV characteristics, as discussed at the beginning of this chapter. According to Eqs. (7.8) and (7.19), the generation rate Gex resulting from this additional photon flux will be: Gex ¼ aðlÞfex ðlÞ. (7.29) The additional free charge carriers generated in this process will lead to an extra current density Jex at wavelength l according to: (7.30) Jex ¼ eð1  RðlÞÞGex ðlÞL Here, L is the effective thickness of the depletion layer and R(l) is the reflection coefficient of the solar cell material at the emission wavelength. Such an approach has been explored in Refs. [32,33].

7.5.3 An example: upconversion in solar cells using Er3þ codoped with Yb3þ Phosphors with trivalent erbium (Er3þ ) that are codoped with ytterbium (Yb3þ ) are a very interesting UC system. The absorption cross section of Yb3þ is 10 times the cross section of Er3þ for parts of the infrared frequency range, which makes it a good sensitizer for Er3þ [41]. Fig. 7.6 shows a model for the resulting coupling scheme. In this scheme, the symbols sif denote the emission and absorption cross sections, and Wif stands for the spontaneous decay rates, where i and f denote the initial and final states, respectively. The ET rates are denoted by Cij , where i and j are the donor and acceptor energy levels, respectively. The dashed arrows indicate instantaneous processes such as MPR. The corresponding rate equations are given by Ref. [41]: dNa ¼  sab fp Na þ sba fp Nb þ Wb Nb dt dNb ¼ sab fp Na  sba fp Nb  Wb Nb þ C2a Na N2 dt Cb0 Nb N0  Cb1 Nb N1  Cb2 Nb N2

(7.31)

(7.32)

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271

(7.33)

dN0 ¼  s02 fp N0 þ s20 fp N2 þ W1 N1 þ C11 N1 N2 þ C22 N22 þ W3 N3 dt þ W4 N4 (7.34) dN1 ¼ W21 N2  W1 N1  Cb1 Nb N1  2C11 N12  s13 fp N1 dt

(7.35)

dN2 ¼ s02 fp N0  s20 fp N2  W2 N2 þ C11 N12  2C22 N22 þ Cb0 Nb Nb N0 dt  C2a Na N2  Cb2 Nb N2 (7.36) dN3 ¼ s13 fp N1  W3 N3 þ Cb1 Nb N1 dt dN4 ¼ s24 fp N2  W4 N4 þ C22 N22 þ Cb2 Nb N2 dt

(7.37) (7.38)

(7.39) NEr ¼ N0 þ N1 þ N2 þ N3 þ N4 In these equations Ni denotes the population density of the energy level i, and fp denotes the photon flux density of the pump wavelength. The values used for the various coupling parameters can be found in Ref. [41]. We further assume ion concentrations of NEr ¼ 2:6  1020 cm3 and NYb ¼ 10  1020 cm3 , and a flux density of about fp ¼ 1:2  1021 m2 s1 at the pump wavelength of 980 nm. Note that this wavelength lies within the absorption spectrum of Si, but we will neglect this adverse process for our simulations. As a result, the current example will become a bit artificial, and strictly speaking it will only be valid for very thin solar cells. Eqs. (7.31)e(7.39) are then solved in the steady state to obtain the nine unknown population densities. For UC in solar cells only the population densities N4 and N3 are of interest. They will emit photons with wavelengths around 550 and 660 nm, respectively. These and other population densities are given in Table 7.2. The photon flux for both N3 and N4 emissions calculated using Eq. (7.3) are found to be: fex ð550 nmÞ ¼ 128  106 cm2 s1 and 6 2 fex ð660 nmÞ ¼ 640  10 cm , assuming a glass layer thickness d ¼ 4000 nm. This photon flux can be used in Eq. (7.30) to calculate the

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Table 7.2 A typical set of population densities, Nx , which were obtained by solving an Er/Yb model system in the steady-state approximation. Parameter Value

Na Nb N0 N1 N2 N3 N4

1:5  1020 cm3 7:5  1021 cm3 2:5  1020 cm3 3:8  1014 cm3 6; 4  1013 cm3 7:6  106 cm3 6:4  106 cm3

generation rate and the resultant photocurrent from the upconverted photons, assuming zero reflection, a depletion layer of size L ¼ 10  105 and absorption coefficients az106 m1 . The additional photon flux can also be used to modify the spectrum of a model photon flux used in solar cell device simulation programs as explained above. For the standard silicon pen junction solar cell that comes prepacked with the WxAMPS package, the efficiency gain from the glass layer is 0.0046%, which is small due to the glass layer thickness used, but good enough to demonstrate the usefulness of our model. Results for the current system are summarized in Table 7.3. Apart from using a very thin conversion layer, it must also be pointed out, that we did not use the most efficient UC system available. In order to properly implement a conversion layer into a standard solar cell device simulation, we would be forced to cater for very thick glass layers compared to the thickness of the pen junction, and thus we will have to employ a multiscale approach [42] to solve Eq. (7.12)e(7.14) on a grid, which is not part of standard solar cell device simulations. Table 7.3 Comparing the basic performance for a standard Si solar cell with and without an additional upconversion (UC) layer (Er/Yb model system). These values were determined using a Si reference system that comes prepacked with the AMPS simulation package, and the effects of the upconversion layer were implemented by modifying a reference solar flux model of the simulation software, according to an upconversion layer of size 0:2 mm. Parameter Without UC With UC

Efficiency, h Open circuit voltage, V oc Short circuit current density, J sc

5.848% 1.6371 V 68.433 A/m2

5.852% 1.6371 V 68.460 A/m2

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7.6 Conclusions We have given an overview of state-of-the-art solar cell device simulations and discussed how to implement augmented light management features like plasmonic nanoparticles and frequency conversion layers. It is interesting to note that important information about the performance of a solar cell can be already be obtained from the Shockley diode equation. Based on the latter, we have pointed out a simple and straightforward way to estimate the generated light currents. This model could be further improved by using simple tunneling and field emission models to estimate the current that effectively passes through the interfaces at the contacts, and by developing simple models to estimate the most important recombination processes due to the trapping and annihilation of free charge carriers. For many experimental groups the use of the Shockley diode equation will be a lot more useful than some of the more detailed device simulations, in particular when researchers are only interested in finding new materials for photovoltaic applications, and in fabricating the simplest types of prototype devices only. We then gave a brief overview of state-of-the-art solar cell device simulations. Here the basic set of equations consists of the Poisson equation for the electric potential across the device, and of two continuity equations, which describe the current densities of the free charge carriers that are generated by the incoming light. These equations have to be supplied with proper physical boundary conditions, and the solar cell itself will be treated as a layered cake model, which may effectively be reduced to a onedimensional numerical problem. This approach covers all of the main processes of light absorption, the corresponding generation and recombination of free charge carriers, and the extraction of photoinduced currents at the contacts. It also allows for light to bounce back and forth between different layers, as happening in a real solar cell. We showed that the key parameters to run these device simulations can all be taken from first principles simulations of bulk materials, which means that for a given material, we can simulate the corresponding photovoltaic device entirely on a computer. Therefore, a solar cell device simulation may start from the atomic structure of the basic materials. Systematic studies along these lines will offer new possibilities to screen large classes of candidate materials for novel types of solar cell devices, in order to avoid many unsuccessful hours in the lab.

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Finally, we also followed up on a promising strategy to develop new type of photovoltaic devices, which is through augmentation with advanced light management features. We discussed how to include these effects into the Shockley diode equation, and into standard solar cell device simulation codes. Our focus was on frequency conversion strategies to harness light outside the spectral range of conventional solar cells, which leads to an additional photon flux and to additional light currents, and we pointed out some of the simplest approaches to implement these processes within the standard theoretical framework of solar cell device simulations. We are well aware of the fact that solar cell research is currently driven by the search for new types of photovoltaic wonder materials like the perovskites [1]. But in the long run, the situation could change quite dramatically, and the focus of the field might switch to a proper “photon management” for just a handful of standard photovoltaic devices. Given the power of the numerical approaches described in this chapter, it could well be that this “light management” approach will primarily be driven by numerical simulations, and not so much by experimental work.

Acknowledgments AQ, IM, and RW would like to thank the Materials for Energy Research Group (MERG) and the DST-NRF Centre of Excellence in Strong Materials (CoE-SM) at the University of the Witwatersrand for support. AQ would also like to thank the Mandelstam Institute for Theoretical Physics (MITP) for support.

References [1] A. Polman, M. Knight, E.C. Garnett, B. Erler, W.C. Sinke, Photovoltaic materials: present efficiencies and future challenges, Science 352 (6283) (2016) aad4424. [2] A. Polman, H.A. Atwater, Photonic design principles for ultrahigh-efficiency photovoltaics, Nat. Mater. 11 (2012) 174e177. [3] S. Fonash, Introduction to Light Trapping in Solar Cell and Photo-Detector Devices, Academic Press, Oxford, 2015. [4] F. Enrichi, A. Quandt, B.C. Righini, Plasmonic enhanced solar cells: summary of possible strategies and recent results, Renew. Sustain. Energy Rev. 82 (3) (2018) 2433e2439. [5] T. Trupke, M.A. Green, P. W€ urfel, Improving solar cell efficiencies by up-conversion of sub-band-gap light, J. Appl. Phys. 92 (7) (2002) 4117e4122. [6] S. Fonash, Solar Cell Device Physics, second ed., Academic Press, Oxford, 2010. [7] J. Nelson, The Physics of Solar Cells, Imperial College Press, London, 2003. [8] N.W. Ashcroft, N.D. Mermin, D. Wei, Solid State Physics, Cengage Learning Asia, Singapore, 2016. [9] W. Shockley, The theory of p-n junctions in semiconductors and p-n junction transistors, The Bell System Technical Journal 28 (1949) 435e489. [10] P. W€ urfel, Physics of Solar Cells, Wiley-VCH, Weinheim, 2009.

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[11] S. Fonash, J. Arch, J. Cuiffi, J. Hou, W. Howland, P. McElheny, A. Moquin, et al., A Manual for AMPS-1D for Windows 95/NT, The Pennsylvania State University 10, USA, 1997. [12] H. Zhu, A.K. Kalkan, J. Hou, S. Fonash, Applications of AMPS-1D for solar cell simulation, AIP Conference Proceedings 462 (1) (1999) 309e314. [13] M. Burgelman, P. Nollet, S. Degrave, Modelling polycrystalline semiconductor solar cells, Thin Solid Films 361 (2000) 527e532. [14] A. Quandt, R. Warmbier, I. Mokgosi, T. Aslan, Solar cell device simulations, transparent optical networks (ICTON), in: 2017 19th International Conference on, IEEE, 2017, pp. 1e5. [15] M.A. Green, Solar Cells: Operating Principles, Technology, and System Applications, Prentice-Hall, Inc., New Jersey, 1982. [16] R.F. Pierret, Semiconductor Device Fundamentals, Addison-Wesley, USA, 1996. [17] A. Luque, S. Hegedus (Eds.), Handbook of Photovoltaic Science and Engineering, John Wiley & Sons, New Jersey, 2011. [18] O. Isabella, A. Smets, K. J€ager, M. Zeman, R. va Swaaij, Solar Energy: The Physics and Engineering of Photovoltaic Conversion, Technologies and Systems, UIT Cambridge Limited, UK, 2016. [19] W. Jones, N. March, Theoretical Solid State Physics: Perfect Lattices in Equilibrium, vol. 1, Courier Corporation, New York, 1973. [20] P.B. Allen, B. Mikovic, Theory of superconductivity Tc, in: H. Ehrenreich, F. Seitz, D. Turnbull (Eds.), Solid State Physics, vol. 37, Academic Press, New York, 1982, pp. 1e92. [21] A. Migdal, Interaction between electrons and lattice vibrations in a normal metal, Sov. Phys. JETP 7 (6) (1958) 996e1001. [22] S. Baroni, S. de Gironcoli, A. Dal Corso, P. Giannozzi, Phonons and related crystal properties from density-functional perturbation theory, Rev. Mod. Phys. 73 (2) (2001) 515e562. [23] S.Y. Savrasov, D.Y. Savrasov, Electron-phonon interactions and related physical properties of metals from linear-response theory, Phys. Rev. B 54 (23) (1996) 16487e16501. [24] F. Giustino, M.L. Cohen, S.G. Louie, Electron-phonon interaction using Wannier functions, Phys. Rev. B 76 (2007) 165108. [25] J. Yan, J.J. Mortensen, K.W. Jacobsen, K.S. Thygesen, Linear density response function in the projector augmented wave method: applications to solids, surfaces, and interfaces, Phys. Rev. B (83) (2011) 245122. [26] I. Tamm, Relativistic interaction of elementary particles, J. Phys.-USSR (9) (1945) 449e460. [27] S. Dancoff, Non-adiabatic meson theory of nuclear forces, Phys. Rev. 78 (4) (1950) 382e385. [28] P. Giannozzi, O. Andreussi, T. Brumme, O. Bunau, M.B. Nardelli, M. Calandra, R. Car, et al., Advanced capabilities for materials modelling with Quantum ESPRESSO, J. Phys. Condens. Matter 29 (46) (2017) 465901. [29] A. Marini, C. Hogan, M. Gr€ uning, D. Varsano, Yambo: an ab initio tool for excited state calculations, Comput. Phys. Commun. 180 (8) (2009) 1392e1403. [30] Y. Liu, Y. Sun, A. Rockett, A new simulation software of solar cells wxamps, Sol. Energy Mater. Sol. Cell. 98 (2012) 124e128. [31] T. Aslan, Simulation of Highly Efficient Solar Cells, University of the Witwatersrand, Johannesburg, South Africa, 2018. Thesis. [32] A. Quandt, T. Aslan, I. Mokgosi, R. Warmbier, M. Ferrari, G. Righini, About the implementation of frequency conversion processes in solar cell device simulations, Micromachines 9 (9) (2018) 435.

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[33] A. Quandt, T. Aslan, I. Mokgosi, R. Warmbier, Advanced light harnessing features in solar cell device simulations, in: 2018 20th International Conference on Transparent Optical Networks (ICTON), IEEE, 2018, pp. 1e4. [34] J. de Wild, J. Rath, A. Meijerink, W. Van Sark, R. Schropp, Enhanced near-infrared response of a-Si: H solar cells with b -NaYF4: Yb3þ(18%), Er3þ (2%) upconversion phosphors, Sol. Energy Mater. Sol. Cell. 94 (12) (2010) 2395e2398. [35] H. Lian, Z. Hou, M. Shang, D. Geng, Y. Zhang, J. Lin, Rare earth ions doped phosphors for improving efficiencies of solar cells, Energy 57 (2013) 270e283. [36] W. Yang, X. Li, D. Chi, H. Zhang, X. Liu, Lanthanide-doped upconversion materials: emerging applications for photovoltaics and photocatalysis, Nanotechnology 25 (48) (2014) 482001. [37] S. Fischer, H. Steinkemper, P. Loper, M. Hermle, J. Goldschmidt, Modeling upconversion of erbium doped microcrystals based on experimentally determined Einstein coefficients, J. Appl. Phys. 111 (1) (2012) 013109. [38] M. Vega, P. Alemany, I. Martin, J. Llanos, Structural properties, Judd-Ofelt calculations, and near infrared to visible photon upconversion in Er 3þ/Yb 3þ doped BaTiO3 phosphors under excitation at 1500 nm, RSC Adv. 7 (17) (2017) 10529e10538. [39] B. Zhou, E. Pun, H. Lin, D. Yang, L. Huang, Judd-Ofelt analysis, frequency upconversion, and infrared photoluminescence of Ho3þ-doped and Ho3þ/Yb3þ-codoped lead bismuth gallate oxide glasses, J. Appl. Phys. 106 (10) (2009) 103e105. [40] B. Walsh, Judd-Ofelt theory: principles and practices, in: Advances in Spectroscopy for Lasers and Sensing, Springer, Germany, 2006, pp. 403e433. [41] L. Shanfeng, M. Zhang, P. Yang, Q. Zhang, Z. Mingshan, Rate equation model analysis on the infrared and upconversion emission of Er/Yb co-doped borate-silicate glass, J. Rare Earths 28 (2) (2010) 237e242. [42] E. Weinan, Principles of Multi-Scale Modelling, Cambridge University Press, Cambridge, 2011.

CHAPTER EIGHT

Light trapping by plasmonic nanoparticles Seweryn Morawiec1, Isodiana Crupi2 1

Institute of Physics, Faculty of Physics, Astronomy and Informatics, Nicolaus Copernicus University, Torun, Poland Engineering Department, University of Palermo, Palermo, Italy

2

8.1 Introduction Scattering of light from subwavelength metallic nanoparticles (NPs) due to localized surface plasmon resonance (LSPR) has been identified over the past decade as a promising approach to realize effective light trapping for thin film solar cells [1e4]. Pioneering work on the plasmonenhanced light coupling into a semiconductor thin film was done in 1998 by Stuart and Hall [5], who demonstrated improved photocurrent generation of a silicon-on-insulator photodetector structure with dense NP arrays placed on top of the device. The applications of plasmonic scattering to photovoltaics appeared only after 2005, with several experimental demonstrations of photocurrent enhancements in silicon solar cells [6e8] supported by theoretical works [9e11]. The resonant interaction of light with the NPs at plasmonic resonance can result in near-field enhancement, far-field scattering, and absorption of light. In particular, the efficient scattering of light over a broad wavelength range from the solar spectrum provides extended path of the light inside the photovoltaic material and consequently increases the probability of scattered photons to be absorbed and generate photocarriers. On the other hand, all the supporting layers additional to the photovoltaic material and all the scattering elements, in particular plasmonic NPs, are the sources of parasitic absorption, meaning that a part of the incident sunlight absorbed by the device does not contribute to the photocurrent [12e14]. As the optical properties of the NPs are strongly dependent on their geometry and on the refractive index of the surrounding mediums, the challenge is to optimize the NPs for (1) maximized scattering and minimized absorption, (2) broad spectral response, (3) broad angular distribution of scattered light, Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00008-2

© 2020 Elsevier Ltd. All rights reserved.

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(4) absorption outside the useful wavelength range, and (5) weak dependence on polarization and angle of incidence. To be directly applicable in the photovoltaic industry, the NP fabrication method needs to be reliable, low cost, and scalable up to square meter sizes. As such, self-assembly processes are most commonly used for the fabrication of the NPs, in particular solid-state dewetting (SSD), consisting in a thermally induced transformation of morphology from a thin film into an array of droplets or NPs. Silver (Ag) is the usual plasmonic material of choice, owing to its high radiative efficiency and low imaginary permittivity in the visible and near-infrared (NIR) spectrum [15]. Plasmonic light trapping with self-assembled NPs can offer a number of advantages: (1) no need for texturing, therefore neither increase of surface recombination nor creation of defects in the photovoltaic material; (2) straightforward scalability; (3) cost-effectiveness, especially when compared with recently proposed light-trapping methods involving nanoimprint lithography; and (4) full compatibility with the current manufacturing technologies in photovoltaics. In the initial designs the NPs were placed on the front surface of the solar cell. Later, it has been demonstrated that a more advantageous arrangement is to couple NPs with the rear mirror in a structure referred as plasmonic back reflector (PBR) [16e19]. In this configuration NPs interact only with the light that was not absorbed during the first pass through the cell material, thus decreasing the optical losses in the visible range [20,21]. This chapter is organized as follows. First, we provide a concise theoretical background on the plasmonic light trapping. Second, we discuss the correlation between the structural and optical properties of selfassembled silver nanostructures, according to the fabrication parameters of the SSD process, on three consecutive stages of integration: NPs on glass substrates, NPs embedded in aluminum-doped zinc oxide (AZO), and NPs assembled in a PBR structure. Third, we discuss the plasmonic enhancement of light absorption in thin silicon films measured by the combination of photothermal deflection spectroscopy (PDS) and Fourier-transform photocurrent spectroscopy (FTPS), a method developed to independently quantify the useful (i.e., generating photocurrent) and parasitic absorption. Last, we discuss the broadband enhancement in thin film hydrogenated amorphous silicon (a-Si:H) solar cells fabricated on optimized PBR structures.

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8.2 Theoretical background Plasmonics exploits the physics of collective oscillations of free electrons in metals, induced by the interactions with electromagnetic radiation. Incident light can be either coupled to the metaledielectric interface in the form of propagating surface plasmon polariton (SPP) waves or confined in the nanoscale metallic structures as localized surface plasmons (LSP). Therefore, plasmonics offers a unique possibility to control optical energy in the subwavelength scale and allows to break the diffraction limit of dielectric photonics.

8.2.1 Localized surface plasmon resonance Localized surface plasmons are nonpropagating excitations of the conduction electrons in a metallic NP placed in an oscillating electromagnetic field, schematically shown in Fig. 8.1A. The interactions with a fixed background of positively charged ion cores give a restoring force on the driven electrons so that a resonance can arise, leading to the strong amplification of electric field both inside and in the near-field zone outside the particle. For noble metals, like silver and gold, the resonance is observed in the visible region of the electromagnetic spectrum. In the first-order approximation, the interaction of an NP with the electromagnetic field can be analyzed by the assumption that its diameter d is much smaller than the illumination wavelength in the surrounding medium, which implies that the phase of harmonic electromagnetic field is nearly constant over the particle volume. As such, the scattering problem can be simplified to a particle in an electrostatic field, known as quasistatic approximation [22]. Let us consider a homogeneous, isotropic sphere of subwavelength radius a located at the origin in a uniform, static electric field, parallel to the z-direction, as depicted in Fig. 8.1B. The surrounding medium is taken to

+ +

--

----

electron cloud

++

+++

metallic particle

--

z

E0

r

T

y ( )

+

+

+++++

(B)

t)

--

E0 e i( kr -

--

(A)

a

y

m

Figure 8.1 (A) Localized surface plasmon resonance mechanism. (B) Geometry used for quasi-static approximation calculations.

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be isotropic and nonabsorbing with dielectric constant εm, and the sphere is described by a complex dielectric function ε. In the quasi-static regime the polarizability, a, of the sphere is given by: ε  εm a ¼ 4pa3 (8.1) ε þ 2εm a is enhanced when the term jε þ2εm j is minimized, which is satisfied by negative values of ε. According to the free-electron model, the real part of the dielectric function for metals is negative for frequencies lower than the plasma frequency, up, which enables (in an oscillating field) a strong resonance of a, known as the dipole mode of LSPR. The radiation of this dipole leads to the scattering of the incident wave by the sphere, and the corresponding cross sections for scattering and absorption, CScatt and CAbs, are given by:

k4 2 8p 4 6

ε  εm

CScatt ¼ jaj ¼ k a
3 6p ε þ 2εm
(8.2)   ε  ε m CAbs ¼ kIm½a ¼ 4pka3 Im ε þ 2εm As CAbs scales with particle volume and CScatt scales with squared volume, the optical response of small particles is dominated by absorption, whereas the response of big particles is dominated by scattering. Although the theory of dipole particles outlined above is strictly valid only for vanishingly small particles, it assures a reasonably good agreement for spheres with dimensions below 100 nm illuminated with visible or NIR radiation. The rigorous solution of Maxwell equations for the case of a spherical particle with arbitrary size can be done with an approach known as the Mie theory [23].

8.2.2 Plasmonic light trapping The unique properties of plasmonic resonances can be exploited for light trapping in thin film solar cells by taking advantage of: 1. Near-field enhancement, schematically shown in Fig. 8.2A. NPs can act as an antenna for the incident light, concentrating the optical energy in the local electric field of the surface plasmon mode. However, for the energy conversion to be efficient, the absorption rate of the semiconductor must exceed the reciprocal of plasmon decay time (typical lifetimes 10e50 fs). This approach is particularly suitable for

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materials wherein small carrier diffusion lengths require ultrathin absorbers for the effective collection of photocarriers, i.e., organic semiconductors, for which an increase in efficiency by a factor of 1.7 has been demonstrated [24]. 2. SPRs, shown in Fig. 8.2B. Incident light can be coupled, e.g., using metallic gratings in rear mirror, into electromagnetic waves traveling along the interface between the metal back contact and the absorber layer. Notably, in this configuration the incident light could effectively be turned by 90 degrees and absorbed along the lateral direction of the solar cell. The penetration depth of evanescent fields into the semiconductor is much larger than that into the metal, thus allowing for the major part of energy to be absorbed in the semiconductor if only its absorption coefficient is sufficiently high to exceed the losses in metal. However, in the visible and NIR spectral ranges, SPPs suffer from relatively high ohmic losses, reducing the propagation lengths and consequently limiting the applicability in solar cells. The corrugations of the mirror can also act as efficient plasmonic scatterers. 3. Far-field scattering, shown in Fig. 8.2C and D. Subwavelength metallic NPs can scatter light efficiently over a broad wavelength range of the solar spectrum due to the LSPR. In addition, when NPs are placed close to the interface between two dielectrics, the scattering occurs preferentially into the material with higher refractive index. Such scattered light is redirected away from the specular direction effectively increasing the path length of light the wavelengths that are normally poorly absorbed in the photovoltaic material. Furthermore, light scattered at angles beyond the critical angle for the total internal reflection (qc of about 16 degrees for the Si/air interface) can be coupled into the guided mode and remain trapped inside the semiconductor slab. The light-trapping scheme relying on far-field scattering has been widely investigated in the past decade and recognized as the most promising for the photovoltaic applications

8.2.3 Optimization guidelines Scattering and absorption of light by plasmonic NPs depend strongly on their size, shape, and surrounding environment. The amount of light scattered (absorbed) by an NP is determined by its CScatt (CAbs), a quantity with dimensions of area given by the ratio of the scattered (absorbed) energy to the incident irradiance [23]. In the literature of single-particle scattering,

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NPs

|E|2 Corrugated mirror

Flat mirror

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A b s o rb e r

(D)

front interface

A b s o rb e r

Tc

A b s o rb e r

A b s o rb e r

(A) Near-field enhancement

Tc

NPs Flat mirror

Figure 8.2 Schematic illustration of plasmon-induced light trapping in thin film solar cells due to (A) near-field enhancement, (B) surface plasmon polariton and plasmonic scattering provided by corrugated back metallic mirror, and (C and D) far-field scattering from metallic nanoparticles placed either on the front (C) or on the rear (D) side of the cell.

however, it is more usual to represent the dimensionless scattering (absorption) efficiency, QScatt (QAbs), which is the ratio of CScatt (CAbs) to the particle cross-sectional area, A: CScatt CAbs (8.3) ; QAbs ¼ A A Extinction efficiency (QExt) is given by the sum of scattering and absorption efficiencies: QScatt ¼

QExt ¼ QScatt þ QAbs (8.4) Certain optimization guidelines, concerning NPs’ geometry and device configuration, can be extracted from theoretical considerations and computational studies present in the literature. Size: Both CScatt and QScatt increase rapidly with the increasing NP diameter up to 100 nm. In addition, the spectral position of LSPR shifts toward longer wavelengths and broadens significantly [11,25]. Therefore, for solar cell applications, big NPs, with sizes of 100e200 nm are generally preferable. On the other hand, bigger NPs may shift the dipolar resonance

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outside of the suitable wavelength range and decrease the scattering efficiency. Shape: The cross-sectional shape of the NP determines their spectral response as well as the efficiency of light coupling into the photovoltaic material, because the light is preferentially scattered toward the material with high refractive index. Hence, the most desirable are geometries with high contact area to the substrate, i.e., cylinder or hemisphere [9]. Environment: High refractive index of the embedding medium is useful for the red shift of LSPR. In addition, when the NPs are placed close to the interface between two dielectric materials, a strong preferential scattering toward the high-refractive-index material occurs. Therefore, for rearlocated NPs, the coupling efficiency increases significantly with decreasing spacing layer between Si and the NPs, which has been demonstrated experimentally [20] and theoretically [10] and can be explained by an overlap between NP near field and the substrate. Device configuration: The high optical losses in the NPs at short wavelengths can be suppressed by placing them on the rear side, so that they interact only with the long-wavelength photons, which are not absorbed during the first pass through the cell absorber material [20,26]. Such configuration has been widely investigated, resulting in experimental demonstration of high photogenerated currents; for instance, 16.9 and 21 mA/cm2 have been achieved as the result of light trapping provided by nanopatterned back reflector (BR) in ultrathin a-Si:H [3] and mc-Si:H solar cells [27], respectively. The nanopatterning is done, however, with complicated and expensive techniques, like electron beam or nanoimprint lithography. An alternative approach takes advantage of self-assembled NPs fabricated in the PBR. It has been demonstrated that PBR can provide efficient light trapping, comparable to state-of-the-art random texturing [17], resulting in photocurrents of 15.6 mA/cm2 achieved with a-Si:H solar cell [28] and 25.5 mA/cm2 with mc-Si:H [29].

8.3 Self-assembled silver nanoparticles 8.3.1 Solid-state dewetting SSD consists in the transformation of morphology from a thin film to an array of islands or droplets due to thermal treatment occurring below the melting temperature of the material. The driving force of the phenomena is the minimization of the free surface energy of the film, the substrate, and the filmesubstrate interface [30,31].

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Thin films are generally metastable in the as-deposited state and would spontaneously dewet to form islands when heated to temperatures well below the melting point of the material, at which surface diffusion is sufficiently high to provide effective mass transport. The threshold temperature at which dewetting is observed decreases with film thickness. The thermodynamics and kinetics of SSD process has been widely studied over the past few decades, first as a degradation mechanism of integrated circuits and microsystems and recently as a large-area manufacturing method of NP arrays with many potential applications. Fabrication of metallic NPs by SSD satisfies the main requirements of photovoltaic industry: simplicity, low-cost, reliability, and direct scalability up to solar panel size.

8.3.2 Nanoparticle fabrication and characterization techniques Ag and AZO films were deposited on soda-lime glass substrates using a radio frequency (RF) magnetron sputtering system, which is fully compatible with the industrial manufacturing lines of thin film solar cells. As-deposited films were found to be continuous, except for the thinnest one (6 nm) having a discontinuous, networked morphology. The thicknesses of the layers were determined from Rutherford backscattering spectrometry (RBS). The postdeposition annealing, in which self-assembled NPs are formed by SSD from the topmost Ag layer, was done at a temperature, TA, in the range between 200 and 500 C in flowing nitrogen. For simplicity, we shall refer to these samples using the abbreviated notation Agx_y, where x is the precursor film thickness, tAg, and y the value of TA. The RBS analyses revealed that no notable out-diffusion of Ag atoms occurs during annealing up to 500 C. The structural properties of the nanostructures were investigated scanning electron microscopy (SEM) using a field emission Zeiss Supra 25 system. The statistical analysis of the NPs was based on the digital analysis of SEM images conducted with the Gatan digital microscope and ImageJ particle analysis software packages. The particles’ area and their in-plane diameters, which for simplicity we should call particle size, were extracted and used to prepare histograms of NP counts and surface coverage (SC) distributions. The optical properties were measured using a Varian Cary 500 doublebeam ultravioletevisibleeNIR spectrophotometer equipped with a 4-inch

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integrating sphere. After the integrated measurements of the total and diffused transmission (TTotal and TDiff, respectively) as well as the total and diffused reflection (RTotal and RDiff, respectively) were performed, the absorption in the sample (Abs) was calculated as: Abs ¼ 100%  TTotal  RTotal  AbsSubstrate

(8.5)

and the light extinction (Ext) was measured as: Ext ¼ Abs þ RTotal þ TDiff  ExtSubstrate

(8.6)

where the background effect of the substrate has been subtracted from original measurements.

8.3.3 Correlation between structural and optical properties of self-assembled nanoparticles The particle ensembles produced by SSD show notably broad distribution of sizes and shapes. Thus the analysis of their morphology and its correlation with optical properties require a statistical approach, commonly realized by histograms of NP counts. However, according to Mie calculations, the absolute values of the extinction cross section for small NPs (w10 nm) are over two orders of magnitude smaller than for those for medium-sized NPs (w100 nm). As such, the simple histogram describing relative counts of NPs does not correlate well with the optical properties of a NP ensemble. A better suited quantity is a histogram of SC, calculated as a sum of the NPs’ in-plane cross-sectional areas in a certain diameter range divided by the total investigated sample area. Such histograms reveal the NP sizes that cover the greatest area and thus have the highest impact on the optical properties of the entire ensemble [25]. Fig. 8.3 shows SEM images, with corresponding size distributions and SC histograms, for tAg of 10.5 nm and three different TA. In the lowannealing-temperature regime, up to 250 C, a formation of irregular NPs was observed and, due to uncertainties in definition of particle size, histograms are not available. Well-defined and ellipsoidal NPs were obtained from TA of 350 C. With a further increase of TA up to 450 C, the size uniformity of the NPs increases, and the number of small particles decreases, which is seen in the narrowing of the particle counts distribution. Furthermore, the mean SC size of NPs, determined as the maximum of Gaussian single-peak fitting to the SC distribution, increases from 136  0.8 up to 149  0.5 nm, with almost unchanged full width at half maximum of 78.5 nm. The total SC decreases from 25.1% for Ag10.5_350

Seweryn Morawiec and Isodiana Crupi

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Figure 8.3 SEM images of Ag NPs formed after annealing of 10.5 nm tAg for1 h at (A) 250, (B) 350, and (C) 450 C. Size and SC distributions of samples Ag10.5_350 and Ag10.5_450 (with well-defined NP size) are shown in (B1, B2) and (C1, C2), respectively. (D) The dependence of the particle area on the particle size at two different annealing temperatures. Points representing perfectly round NPs lie along the pr2 curve. NPs, nanoparticles; SC, surface coverage; SEM, scanning electron microscopic. Reprinted from S. Morawiec, M.J. Mendes, F. Priolo, I. Crupi, Plasmonic nanostructures for light trapping in thin-film solar cells, Mater. Sci. Semicond. Process. 92 (2019) 10e18; with permission of Elsevier.

to 23.5% for Ag10.5_450, and since no noticeable loss of silver occurred during the annealing process, the mean height of the NPs is expected to increase with TA. Furthermore, the shape uniformity in terms of in-plane roundness of the particles increases with annealing temperature, as shown in the plot of particle area against particle size (longitudinal diameter) in Fig. 8.3D. NPs smaller than 100 nm are approximately round for both Ag10.5_350 and Ag10.5_450, whereas the bigger NPs tend to be more elongated, yet they are rounder when formed at 450 C.

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A similar evolution of NPs’ morphology was also observed for tAg of 7.5 nm, annealed in the same range of temperatures. For thickness of 16 nm, a TA of 200 C was found to be the threshold value for nanostructure formation, below which the morphology of the film does not change significantly with respect to the as-deposited film. Also, an increase of TA up to 400 C leads to the formation of morphology with a large number of small NPs spread among a fewer number of much larger particles resulting in a double-peak distribution of the SC histogram. Further increase of TA up to 500 C causes a rapid decline of the number of small NPs and a transition to a single-peak SC distribution. The observed mechanism of nanostructure evolution can be explained in terms of the Ostwald ripening process. For sufficiently high TA, the surface diffusion length is higher than the average distance between the particles, allowing their aggregation and the formation of more energetically favored big particles. High atomic diffusion enables also the formation of uniform round-shaped NPs, because it tends to minimize the particles’ surface tension. The structural analysis of more than 30 samples evidenced that the obtained nanostructures can be categorized according to their morphology into four groups: thin films (F), irregular NPs (I), double-peak (D), and single-peak (S) distributions. Each group is formed in a specific range of fabrication parameters that can be represented as a certain area on the tAg against TA diagram shown in Fig. 8.4. Such graph can be considered as a structural-phase diagram of nanostructure formation. The temperature threshold for dewetting increases with film thickness, in agreement with the dependence found by Kim et al. [32]. For layers thinner than 12 nm, a direct transition from I-type to S-type nanostructure with increasing TA was observed. Above 12 nm, D-type nanostructures are formed between I-type and S-type regions, whose domain widens significantly with increasing tAg. It was also observed that extended annealing time does not change the temperature threshold either for film percolation or for the transition from I- to S- or D-type nanostructure. Fig. 8.5 shows the mean size of the SC distributions as a function of tAg. When the NPs are characterized by a D-type distribution, the mean SC size was shown for both the small and large particle peaks. An almost linear increase in the mean SC size with the precursor film thickness was observed. This demonstrates that tAg is the main parameter determining the mean SC size, whereas TA plays major role in NPs’ shape. A linear dependence of average particle size against precursor film thickness for dewetted gold films has also been obtained by Kojima et al. [33].

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Figure 8.4 Structural-phase diagram of NPs formation. The morphology of annealed thin films can be categorized into thin film (F-type, violet squares), irregular NPs (Itype, blue diamonds), double-peak surface coverage size distribution (D-type, green triangles), and single-peak surface coverage size distribution (S-type, red circles). The background SEM images illustrate each type of morphology. NPs, nanoparticles; SEM, scanning electron microscopic. Reprinted from S. Morawiec, M.J. Mendes, S. Mirabella, F. Simone, F. Priolo, I. Crupi, Self-assembled silver nanoparticles for plasmon-enhanced solar cell back reflectors: correlation between structural and optical properties, Nanotechnology 24 (26) (2013) 265601; with permission of IOP Publishing.

The optical properties of the nanostructures obtained from three precursor films, after annealing at different TA, are shown in Fig. 8.6. All investigated I-type nanostructures (samples Ag7.5_250, Ag10.5_250, and Ag16_250) exhibit wide spectral response as expected from the large spread in NP size and shape. As the NPs become bigger with increasing tAg, the RTotal and Abs spectra broaden significantly. The absorption of the NPs in the sample Ag7.5_250 clearly dominates over their scattering properties in the entire wavelength range, whereas for Ag10.5_250 and Ag16_250 the reflection becomes higher than absorption for wavelengths above 500 nm. For particles formed from films below 12 nm, the increase in TA leads to a clear narrowing of the spectral response, which is related to the formation of the S-type structure with improved homogeneity of sizes and shapes. For both thicknesses (7.5 and 10.5 nm), a comparison between samples formed at 350 C and 450 C shows a clear decrease of absorption and increase of total reflection and scattering for higher TA. This results from the increased mean size of the NPs as well as from the reduction in the number of small particles. On the other hand, the spectral narrowing of the optical

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Figure 8.5 Dependence of the mean SC size on film thickness for well-defined Ag NPs (S and D-type). The solid symbols indicate the mean SC size of the small particles’ peak in D-type distribution. NPs, nanoparticles; SC, surface coverage. Adapted from S. Morawiec, M.J. Mendes, S. Mirabella, F. Simone, F. Priolo, I. Crupi, Self-assembled silver nanoparticles for plasmon-enhanced solar cell back reflectors: correlation between structural and optical properties, Nanotechnology 24 (26) (2013) 265601; with permission from IOP Publishing.

characteristics results from narrower size distribution and increased roundness of the NPs. Two distinguishable peaks are present in the spectra: a narrow peak at w360 nm and a broad one at longer wavelengths. The latter corresponds to the dipolar mode of plasmonic resonance, which red-shifts with increasing NP size (for RTotal from 450 nm for Ag7.5_450 to 510 nm for Ag10.5_450). On the other hand, the peak remains at approximately the same wavelengths with increasing TA because of two contradictory effects: an increase in the particle size, which red-shifts the resonance, and a reduction in the elongation of the particles’ shape and an increase of its height making it more spherical, which blue-shifts the resonance. The shortwavelength peak can originate from different effects: a plasmonic resonance of the small NPs (dominated by absorption), a quadrupole mode sustained by the bigger NPs, and a mode related to the shorter diameter of ellipsoidal NPs. The nanostructures fabricated from tAg of 16 nm have a wide optical response. The increase of TA from 250 C to 500 C leads to a reduction in the parasitic absorption with only a slight decrease in diffused reflection and transmission. The decrease in RTotal with increasing TA can be, however, related to the decrease in the total SC from 31.5% for Ag16_250 to 19.5% for

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Figure 8.6 Optical properties of 7.5 nm (left column; plots A, D and G), 10.5 nm (center column; plots B, E and H), and 16 nm (right column; plots C, F and I) Ag film annealed in the range between 250 and 500 C. Adapted from S. Morawiec, M.J. Mendes, S. Mirabella, F. Simone, F. Priolo, I. Crupi, Self-assembled silver nanoparticles for plasmon-enhanced solar cell back reflectors: correlation between structural and optical properties, Nanotechnology 24(26) (2013) 265601; with permission from IOP Publishing.

Ag16_500. These structures exhibit a less clear distinction between the resonance peaks because of the significant broadening of the optical response with increasing NP size, as well as for the additional higher-order modes, which reduce the resonant character plasmonic excitations. Fig. 8.7A depicts the spectral position of the dipolar resonance wavelength, in terms of the extinction peak, with increasing mean SC size of the NPs. An approximately linear correlation was found to hold, apart from samples annealed at 300 C, which have resonance at shorter wavelengths than samples annealed at higher temperatures. This can be due to different mean height of the NPs. By increasing the mean SC size from 80 to 220 nm the resonance shifts from 420 to 580 nm, which gives an average shift of 1.05  0.07 nm/nm determined from a linear fit for samples annealed in the 350e500 C range. The linear relation reveals that the main parameter determining the resonance wavelength is the size of the NPs, which is directly proportional to tAg (see Fig. 8.5). The most important issue for application in solar cells is the competition between optical losses and scattering properties of the NPs. Scattering is beneficial for generation of photocarriers, whereas parasitic absorption in the NPs leads to a drop in the cell efficiency. An important quantity is

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Mean SC size (nm) Figure 8.7 (A) Dependence of the resonance wavelength on the mean SC size and (B) correlation between the integrated scattering performance and the mean SC size of well-defined Ag NPs (S-type and D-type). The inset shows a TEM image of NPs formed from 12 nm thick Ag film annealed at 400 C for 1 h. NPs, nanoparticles; SC, surface coverage; TEM, transmission electron microscopic. Adapted from from S. Morawiec, M.J. Mendes, F. Priolo, I. Crupi, Plasmonic nanostructures for light trapping in thin-film solar cells, Mater. Sci. Semicond. Process. 92 (2019) 10e18; with permission of Elsevier and S. Morawiec, M.J. Mendes, S. Mirabella, F. Simone, F. Priolo, I. Crupi, Self-assembled silver nanoparticles for plasmon-enhanced solar cell back reflectors: correlation between structural and optical properties, Nanotechnology 24 (26) (2013) 265601; with permission from IOP Publishing.

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the integrated scattering performance (Gscatt), over the whole measured wavelength range, defined as: R R TDiff þ RDiff R R R Gscatt ¼ (8.7) Abs þ TDiff þ RDiff The Gscatt, plotted in Fig. 8.7B, shows a rapid increase with increasing mean SC size and saturation behavior above w160 nm. However, whereas the theoretical scattering efficiency is expected to saturate at values above 90% [25], the experimental Gscatt saturates at around 50%. Such significant discrepancy between theoretical and experimental study is mostly due to high parasitic absorption of the NPs, which can be attributed to the polycrystalline nature and defects in the material, as revealed by the transmission electron microscopic image shown in the inset of Fig. 8.7B. In this case, collectively oscillating electrons dissipate their energy in the interactions with grain boundaries, defects, or surface roughness [34,35]. Other sources of absorption can be the surface roughness and the sharp corners, which act as hot spots (high enhancement of the near field) dissipating the plasmon energy. An important issue affecting the optical properties of NPs is the tarnishing of silver when stored in air. Silver is well known to react with hydrogen sulfide (H2S), present in the atmosphere due to the air pollution, to form a metallic silver core- and silver sulfide shell structure (Ag0/Ag2S) [36]. Because Ag2S is a strongly absorbing material in the visible range, the sulfidation of the NPs results in a considerable decrease of the plasmonic resonance strength. The aging effect is depicted in Fig. 8.8, referring to NPs formed from 12 nm thick Ag film annealed at 400 C, by comparing their optical properties immediately after fabrication and after 8 weeks’ storage in air atmosphere. In addition, a slight 10 nm red-shift of the Ext peak was observed after storage. AZO is an important material commonly used in thin film silicon solar cell technology as a transparent contact layer placed between n-type Si and the metallic mirror [37,38]. This provides a barrier for diffusion of metal atoms into the absorber [39], improves band alignment, and lowers parasitic loses in the mirror when compared with the metal deposited directly on silicon [40]. As discussed in the introduction chapter, the NPs would be preferably located at the rear side of the cell and embedded in the AZO layer between Si and BR. An encapsulation of the NPs in zinc oxide leads also to their improved stability against the sulfidation occurring in atmospheric air. It is therefore important to study the properties of self-assembled Ag NPs formed on and embedded in the AZO.

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Figure 8.8 The effect of Ag tarnishing on the optical properties of the nanoparticles fabricated from tAg of 12 nm annealed at 400 C for 1 h. Solid lines correspond to the measurements done immediately after fabrication, and dashed lines correspond to the measurements after 8 weeks in air.

Fig. 8.9 depicts the SEM images together with the corresponding size and SC distributions of NPs, formed from 12 nm thick Ag film annealed at 400 C, fabricated in the exact same process on (Fig. 8.9A) bare glass and (Fig. 8.9B) glass coated with 50 nm AZO. NPs formed on AZO tend to be larger than those formed on glass, although the maximum sizes of NPs are approximately the same in both cases. The average NPs’ size and the mean SC size, determined from Gaussian-type peak fitting, increase from 139  2 nm to 169  2 nm and from 164  2 nm to 190  1 nm, respectively. The relative number of small particles and the overall number of particles are reduced on AZO. In addition, the total SC increases from 20.3% to 26.1%, which reveals that NPs formed on AZO also tend to be flatter than those formed on glass substrate. All the morphological differences described above can be attributed to the higher surface free energy of AZO than glass. Besides the important morphological changes, formation of the NPs on AZO-coated glass also influences their optical properties due to the increase of the effective refractive index of the medium surrounding the NPs. Fig. 8.10 shows the comparison between the Ext spectra of particles formed in the same process on (1) bare glass, (2) glass coated with 50 nm thick AZO, and (3) on 50 nm thick AZO coated with another 50 nm thick AZO layer. The Ext peak was found to red-shift from 555 nm for NPs on glass, to

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Figure 8.9 SEM images of Ag NPs formed on (A) bare glass and (B) glass coated with 50 nm of AZO after annealing of 12 nm thick precursor film for 1 h at 400 C. The corresponding NPs’ size and SC distributions are shown in (A1, B1) and (A2, B2), respectively. NPs, nanoparticles; SC, surface coverage; SEM, scanning electron microscopic. Reprinted from S. Morawiec, M.J. Mendes, F. Priolo, I. Crupi, Plasmonic nanostructures for light trapping in thin-film solar cells, Mater. Sci. Semicond. Process. 92 (2019) 10e18; with permission of Elsevier.

643 nm for NPs on AZO, and to 733 nm for NPs embedded in AZO. Furthermore, a slight rise of the peak intensity and a significant peak broadening were observed. The differences between the first two cases can originate from both the increased refractive index and the increased mean NP size. However, the differences between the last two are purely an effect of a higher refractive index. Both effects of the spectral red-shift and broadening are advantageous for application in photovoltaic, due to the better spectral matching of the plasmonic resonance with the wavelengths that are normally poorly absorbed in silicon. As the contribution of the substrate has not been subtracted from the measurements shown in Fig. 8.10, absorption edges of glass and AZO are clearly observed at 350 and 400 nm, respectively. Moreover, the out-of-resonance values of Ext approach the values of the total reflection of the substrate. The optical properties of the NPs were measured in two distinct configurations: front-side illumination, when the incident light impinges the NPs first, and rear-side illumination, when the incident light impinges the glass substrate first. The two configurations correspond to the way in which

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Figure 8.10 Extinction measurements of nanoparticles formed from 12 nm thick Ag film annealed at 400 C for 1 h in three different configurations. Solid lines correspond to the front-, and dashed, to the rear-side illumination.

NPs interact with the sunlight in the substrate and superstrate thin film solar cells, respectively. The Ext spectra for front- and rear-side illumination, shown in Fig. 8.10 with solid and dashed lines, respectively, were found to remain almost identical in the whole investigated spectral range. This reveals that the resonance strength is equal in both cases. Nevertheless, significant differences were observed for the single components of the extinction, depicted in Fig. 8.11 for the three samples described previously. The main drawback of rear-side illumination is a clearly lower RTotal, which for constant Ext leads to a notably higher Abs. On the other hand, the TDiff is slightly increased and RDiff is almost doubled, whereas their spectral shape is moderately narrower. The asymmetry between the two illumination configurations should be attributed to the asymmetry of the refractive index around NPs as well as to their hemispherical cross-sectional shape. Apparently, when illuminated from the rear side the sharp corners of the NPs act more pronouncedly as hot spots, increasing the resonant interactions, which gives rise to both absorption and scattering. The effects described above stay in good agreement with the results of Pennanen and Toppari [41], who investigated the efficiency of light coupling into a planar waveguide due to gold nanodisks and found that forward scattering is approximately equal regardless of the illumination direction, whereas the backscattering is much stronger toward the high-refractive-index material.

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Figure 8.11 Optical properties of NPs formed from tAg of 12 nm annealed at 400 C for 1 h in different configurations: NPs on glass (A and D), NPs on 50 nm thick AZO (B and E), and NPs embedded in 100 nm thick AZO (C and F). Left column corresponds to the front and right column to the rear-side illumination. NPs, nanoparticles. Reprinted from S. Morawiec, M.J. Mendes, F. Priolo, I. Crupi, Plasmonic nanostructures for light trapping in thin-film solar cells, Mater. Sci. Semicond. Process. 92 (2019) 10e18; with permission of Elsevier.

8.3.4 Plasmonic back reflectors Parasitic absorption in metallic nanostructures can be a major obstacle for achieving efficient light trapping in thin film solar cells if the benefit coming from light scattering cannot exceed losses in the nanostructures. Thus, it is preferable to incorporate the NPs in the BF of the solar cells so that they only interact with the wavelengths that are not fully absorbed in a single pass through the cell. A typical configuration of a PBR consists of flat silver mirror, a thin layer of a transparent conductive oxide (TCO) spacer, metallic NPs, and TCO cover layer. The main optical quantity that describes the suitability of a PBR for light trapping is RDiff as it determines the amount of light scattered away from the specular direction and, thus, the amount of photons that have increased probability of being absorbed due to the longer path in the solar cell and contribute to the absolute increase of the external quantum efficiency

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(EQE). The total reflection indicates how close a particular PBR is to the perfect mirror and thus determines the parasitic losses with the relation: Abs ¼ 100%  RTotal (8.8) The contribution of the PBRs’ specular reflection (RTotal - RDiff) is only significant at the wavelengths that are not fully absorbed during the first pass through the cell but have high probability of being absorbed in the second pass after the specular reflection. In the near bandgap region, due to the rapid decrease of the absorption, RDiff plays a major role because the path length required for absorption is much longer than the cell thickness. Therefore, a first-order estimation of PBR performance can be done by considering the values of RDiff in a light-trapping range suitable for specific material. For typical a-Si:H solar cells with 300 nm thick intrinsic region, the wavelength range in which light trapping can provide significant photocurrent gain extends from 600 to 800 nm, which is the range between the full extinction of light in a single pass through the cell and the bandgap of a-Si:H. The morphology of the NPs has a strong impact on the optical properties of the PBRs and can be tuned by the SSD parameters. It is generally desirable to prevent the formation of small or irregularly shaped NPs, as they are known to increase the parasitic absorption. At the same time, big NP (greater than 200 nm) are not desirable as they can degrade the electrical properties of the cells due to the roughness induced in the Si layers [28,42]. The impact of the AZO spacer layer thickness on the scattering and parasitic absorption was investigated by Sesuraj et al. [43] who reported an optimum value in the range of 40e50 nm, depending on the size of the NPs. Tan et al. [28] reported that an AZO cover layer of at least 60 nm flattens out the sharp surface features caused by the NPs, preventing them from deteriorating the quality of the a-Si:H cell layers. PBRs were fabricated on soda-lime glass substrates by sequential deposition of a 100 nm thick Ag mirror, a 40 nm thick AZO spacer layer, and a thin precursor Ag film, using RF magnetron sputtering. The deposition was followed by annealing in flowing nitrogen, in which self-assembled NPs are formed by SSD from the topmost thin Ag layer. To embed the NPs, a coating of 80 nm AZO was deposited as the last fabrication step. Fig. 8.12 depicts the comparison between NPs formed from 12 nm Ag film after annealing at 200 and 400 C for 1 h. First, samples fabricated at lower temperature exhibit significantly worse performance in both RTotal and RDiff. It originates from the irregularly shaped NPs and high SC obtained at low annealing temperatures. As shown in the SEM in Fig. 8.12C and D, a

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Figure 8.12 (A) Total and (B) diffuse reflection of the PBRs with NPs formed from tAg of 12 nm annealed at 200 C and 400 C for 1 h. Corresponding NPs’ morphologies are shown in the SEM images (D and C), respectively. NPs, nanoparticles; PBR, plasmonic back reflector; SEM, scanning electron microscopic.

pronounced change in nanostructure morphology from elongated islands to round NPs is observed with the increase in temperature from 200 to 400 C. In addition, the reduction in the number of small particles and decline of the SC is observed. This reveals the requirement of the elevated temperatures for NP formation, in accordance with the previous discussion. Such high temperatures forbid the implementation of SSD-formed NPs as a postprocess on top of thin film silicon cells, because the SieH bonds are damaged at temperatures w200 C. Therefore, the main advantage of PBR configuration is that the a-Si:H p-i-n structure is not affected by the annealing step performed before the deposition of solar cell. Fig. 8.13 depicts the annealing time-dependent evolution of the NPs’ morphology formed from 8 nm Ag film annealed at 400 C for 0.5, 1, and 4 h and the corresponding changes of their optical response. A similar evolution was also observed for NPs’ formed from tAg of 12 nm [21]. The extended annealing time clearly favors the formation of more uniform

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distributions of sizes and shapes of NPs resulting in a desirable higher RTotal and, what is more important, also higher RDiff of the PBR and a blue-shift of the resonance peak. The samples annealed for 0.5 and 1 h contain similar number of small particles, and thus both have comparable RTotal in the 400 to 600 nm range. In contrary, the small NPs are almost not present in the sample with annealing time extended to 4 h, resulting in lower parasitic absorption and hence higher RTotal. The lower parasitic absorption in redNIR region can be attributed to more regular NP shape, lower SC, and possible improvement of NPs’ crystallinity. At the LSPR frequencies, the ratio of RDiff to RTotal, often termed “haze,” exceeds 90%. This proves that almost all incident photons interact with the NPs because all the light that is not absorbed is reflected diffusively from the PBRs. Therefore, the main factor limiting the performance of the PBRs is the parasitic absorption. In this respect, there is a significant discrepancy between theoretical calculations and experimental results because,

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according to single-particle electromagnetic modeling, Ag NPs with sizes above 100 nm should exhibit dominant scattering and negligible absorption cross section [9,11,43,44], as discussed in Section 8.2.3. Although the origin of the unexpectedly high absorption remains unclear, it can be attributed to dissipative interactions between the NPs, between the NPs and the mirror, the roughness and sharp features of the NPs, and their polycrystalline nature.

8.4 Plasmon-enhanced absorption in thin silicon films 8.4.1 Independent quantification of useful and parasitic absorption The coupling of plasmonic particles with photovoltaic material results in the enhancement of light absorption. The important problem in plasmoninduced light trapping is to discern between the absorption originating in the photovoltaic material, i.e., useful absorption, and absorption originating in all supporting layers including the metallic NPs, i.e., parasitic absorption. The trade-off between the beneficial effects of scattering and the deteriorating effects of parasitic absorption can severely limit the overall photocurrent enhancement that can be produced in solar cells. The contributions of useful and parasitic absorption are inseparable with standard transmission-reflection spectroscopy because the optical properties of metallic NPs depend strongly on the dielectric function of the surrounding material. Therefore the contribution of metallic NPs cannot be measured separately and simply subtracted from the measurement of a complete sample comprising both NPs and photovoltaic absorber. Nonetheless, such measurement is possible by a combination of two absorption spectroscopy techniques [14]: namely photothermal deflection spectroscopy (PDS) and Fourier-transform photocurrent spectroscopy (FTPS). In this methodology, PDS accounts for all absorption that results in generation of heat. On the contrary, the FTPS is derived from the number of photocarriers generated in the photovoltaic absorber and collected on external electrodes using applied bias. Therefore, it is proportional to the useful absorption. Both techniques exhibit superior sensitivity over transmittance/reflectance spectrophotometry and have been extensively used to analyze electronic defects in semiconductors. The difference between PDS and FTPS determines the total parasitic absorption, which receives a contribution from all supporting layers and scattering elements. As such, the combination of PDS and FTPS performed sequentially on the same sample allows for the independent quantification of the useful and total parasitic absorption, as attained

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by the conventional 1-R and EQE measurement on the complete solar cell structure. However, it can be performed without the need to process the complete device. For the PDS, a Fluorinert electronic liquid FC-72 with a low refraction index of 1.25 was used to simulate conditions similar to ambient atmosphere. To collect the FTPS signal, a specific arrangement was used comprising a coplanar front electrode made of an electrolyte and a transparent conductive oxide sandwich structure [45].

8.4.2 Absorption enhancement in mm-Si thin films With the aim of investigating the light trapping in substrate configuration thin film solar cells, self-assembled silver NPs were incorporated in distinct arrangements, considered as different stages of completion of a PBR. The NPs were fabricated by SSD of 12 nm thick Ag films annealed at 400 C for 1 h in nitrogen atmosphere on (1) 50 nm thick AZO-coated glass (SAZO_NPs) and (2) a stack of 100 nm thick flat Ag BF and 50 nm thick AZO spacer layer (SBR_NPs). The reference samples without NPs (SAZO and SBR, respectively) were fabricated in the same processes (see schemas in Fig. 8.14B to D). According to the discussion presented in Section 8.2.3, the optical characterization (not shown here) of substrates before the deposition of Si films revealed that the reference samples without NPs show negligible diffused reflectance, which increases in the presence of NPs. A substantial enhancement is observed when the NPs are coupled with a flat BF, because the mirror not only reflects the transmitted diffuse light from the NPs but also gives rise to a constructive near-field interaction. Subsequently, the samples were used as substrates for the deposition of 0.9 mm thick mc-Si:H by plasma-enhanced chemical vapor deposition (PECVD). The deposition conditions assure high degree of crystallinity required for the narrow bandgap, which allows for the optical absorption to extend up to 1100 nm in the NIR. On the other hand, full extinction of light in a flat film was measured for wavelengths shorter than 500 nm. Thus, light trapping can provide absorption enhancement within the 500 to 1100 nm window. The PDS and FTPS absorption spectra of 0.9 mm mc-Si:H deposited on samples SAZO, SAZO_NPs, and SBR_NPs are plotted in Fig. 8.14. The only sources of optical losses in the reference sample are the thin AZO layer and the glass substrate, therefore both PDS and FTPS signals overlap,

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Figure 8.14 (A) Total (PDS) and useful (FTPS) absorption spectra of 0.9 mm thick mc-Si:H deposited on SAZO (B), SAZO_NPs (C), and SBR_NPs (D). The FTPS absorption for sample SBR_NPs with additional 50 nm thick AZO layer separating NPs and silicon is shown for comparison (SBR_NPsþAZO). The marked area represents the parasitic absorption for SBR_NPs being the difference between PDS and FTPS spectra. The inset illustrates schematically light trapped in Si by total internal reflection. FTPS, Fourier-transform photocurrent spectroscopy; PDS, photothermal deflection spectroscopy. Reproduced , A. Vetushka, M. from S. Morawiec, J. Holovský, M.J. Mendes, M. M€ uller, K. Ganzerova Ledinský, F. Priolo, A. Fejfar, I. Crupi, Experimental quantification of useful and parasitic absorption of light in plasmon-enhanced thin silicon films for solar cells application, Sci. Rep. 6 (2016) 22481; under Creative Commons license.

indicating that practically all absorption occurs in silicon. It should be noted that the relatively low absolute values of absorption below 80% originate from the lack of the transparent contact deposited on top of silicon, which plays an important role as an antireflection coating. The deposition of mc-Si:H on the substrates containing the NPs resulted in considerable enhancement of useful absorption in the entire investigated wavelength range. The parasitic losses, calculated as the difference between PDS and FTPS spectra, start to play a role only for the wavelengths above 730 nm, which are already poorly absorbed in the flat film, and increase significantly toward the bandgap of mc-Si:H. This behavior can be explained by assuming that the total parasitic absorption arises from multiple interactions between the incident light and the NPs. The number of such interactions for weakly absorbed NIR light trapped in a silicon slab (depicted schematically in the inset of Fig. 8.14) increases substantially with decreasing absorption coefficient, and hence increasing illumination wavelength. Consequently, the vast number of interactions results in significant overall

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losses even if, as predicted by theoretical calculations [10,11], the absorption in a single interaction is small compared with scattering. Therefore, in the 500 to 700 nm wavelength range, where the light has a high probability to generate photocarriers after only a few scattering events, parasitic absorption plays only minor role. The scattering properties of particles are strongly dependent on the dielectric function of the embedding medium, and therefore also on the distance between the NPs and the high-refractive-index material [10,20]. The FTPS signal of sample SBR_NPs with additional 50 nm thick AZO layer separating NPs and silicon (plotted in Fig. 8.14) exhibits a significant blue-shift of the absorption edge and thus a clearly lower enhancement of useful absorption in the NIR region. This effect can be attributed to the narrowing of the angular distribution function of scattered light, hence lower light coupling efficiency, with decreasing overlap between the NPs’ near field and the silicon. The spacer layer in solar cells is required for a barrier for the diffusion of metal atoms into the silicon, thus preventing from the deterioration of the electrical properties of the doped layer and consequent degradation of the p-i-n junction. Nonetheless, the separation should be kept as thin as possible to ensure high enhancements of useful absorption. On the other hand, a thicker AZO layer flattens out sharp features and improves the morphology of the surface, which is beneficial for the growth of dense mc-Si:H films with lower density of defects and better electrical performance [43]. The main sources of parasitic losses can be attributed to the dissipative interactions of light with the NPs, the Ag mirror, and the AZO spacer layer. It should be noted that, even though no significant parasitic absorption is observed in Fig. 8.14 for the sample with only the AZO layer (SAZO), the losses in this layer or in the flat Ag mirror will necessarily increase with the presence of the nearby NPs. The conformal growth of mc-Si:H films on top of the NPs results in a random texturing of their front surface. The surface morphologies were found to be fairly similar for the films deposited on substrates containing NPs having a root mean square (RMS) roughness ranging from 20 to 24 nm. As such, two distinct light-trapping mechanisms are contributing to the overall useful absorption: (1) improved light coupling into silicon (antireflection action) provided by the surface texture and (2) scattering of light by the plasmonic NPs. The high-energy photons are entirely absorbed in their first pass through the mc-Si:H film and therefore do not reach the back side of the film and have no possibility to interact with the NPs. As such, the higher absorption measured at wavelengths shorter than 500 nm

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for samples with NPs can only be attributed to the front surface texture. In the light-trapping window, both mechanisms can give rise to the observed enhancement of useful absorption and their contributions cannot be measured separately. Nonetheless, computational studies [46,47] indicate that small RMS roughness ( 22.1% in 2017 [64,65]. Despite impressively high efficiencies, PSCs face critical challenges, such as reduced environmental stability which prevents them from competing with established technologies; and there is still room for tailoring charge recombination, both in the perovskite and at the device interfaces, to increase performance. For instance, the carriers’ diffusion length in state-ofthe-art PSCs based on methylammonium lead iodide (MAPbI3) is about 100 nm, due to its poly-crystalline film structure, therefore devices with an ultrathin perovskite absorber (below the typical 500-nm thickness used) could be beneficial for reducing recombination losses. As such, optical solutions are promising avenues for improving not only the PSCs efficiency, by allowing physically thinner but optically thicker devices, but also their market applicability by enabling higher mechanical flexibility due to the reduced thickness (in the same way as targeted for a-Si:H technologydsee Section 9.3.1.2). Recently, several LT solutions have been investigated for PSCs [52], for instance, by implementing ultrathin textures on the perovskite [66,67] or the substrate [51,68], applying ray optic schemes by prism arrays [69], microlens [70], or nanocone [71] arrays, plasmonic

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metal nanoparticles [72e74], light grating-coupled surface plasmon resonances [75], and embedding TiO2 nanoparticles into the mesoporous TiO2 layer [76]. However, most of these approaches require nano-/microstructuring the perovskite layers, thereby imposing the same disadvantage of conventional textured-based LT due to the severe compromise between optical benefits and electrical deterioration. This can be a particularly crucial drawback for PV devices whose carrier lifetime is chiefly limited by interfacial recombination/trapping, as has been shown to be the case for PSCs [13]. Besides, none of these LT approaches has yet shown performance enhancements superior to those presented in Section 9.3.1 for thin-film silicon cells. Higher PSC improvements are envisaged with the front-located wavelength-sized LT structures described in the previous section. In particular, the TiO2 void arrays are highly promising for application in PSCs with an inverted (substrate-type) cell architecture (see Fig. 9.5), in which the light enters through the cell side and not from the transparent substrate as in conventional superstrate PSCs. In addition to the superior broadband absorption gains that such wave-optical structure can provide (e.g., relative to the domes geometry), its integration in PSCs has important advantages besides the optical ones [46]: (1) Material compatibilitydTiO2 is the typical electron transport material (ETM) used in the n-type contact of state-of-the-art PSCs, so the LT structure can be realized by a photonic-structured ETM which is conformally coated with a TCO layer made of ITO acting as the front electrode (see Fig. 9.5C). (2) Electrical preservationdThe integration of the photonic features in the solar cells’ ETM, above the perovskite absorber, enables the implementation of LT without structuring the perovskite layer, i.e., the roughness and consequent recombination (due to induced defect density) in the cell is not increased since the perovskite remains planar (in contrast with conventional texturing approaches). This is particularly beneficial for PSCs, since their main recombination mechanism is caused by surface defects [13]; and due to the fact that they are usually deposited by spin-coating techniques, so the uniformity and electronic properties of the patterned layers are substantially improved when coated on flat surfaces. (3) UV blocking for improved stabilitydThe front-coated TiO2 LT structures revealed to act as highly effective UV-blocking layers, able to protect the underneath sensitive layers from UV exposure which is one of the main causes of degradation of PSC technology [77,78], while

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improving light absorption at longer wavelengths and, thus, increasing the cell efficiency. (4) Broader applicabilitydThe photocurrent enhancement caused by the LT features allows reducing the perovskite absorber thickness (below the conventional w500 nm), which enables improved flexibility and cost reduction [52,79]. Besides, the LT integration in an inverted PSC architecture enables the use of a much wider range of substrates, for instance, with flexible opaque materials (e.g., PI, metal foils, fabrics [80], plastics [81], paper-based [5]), relative to the conventional superstrate architecture which can only be supported on transparent substrates. In addition, the inverted PSC configuration can be directly integrated onto multijunction stacks to realize higher-efficient tandem devices (e.g., wide-band gap perovskite top cell coupled with Si or CIGS bottom cell) [82e84]. Fig. 9.5 and Table 9.1 present the results attained with the optimized hexagonal array of semispheroidal voids integrated in the TiO2 ETM of PSCs with distinct perovskite thickness: 250 nm (ultra-thin) and 500 nm (conventional). The absorption spectra (red curves), given in Fig. 9.5A and B, respectively, are compared with those of the unstructured reference PSCs (blue curves) with an optically optimized set of thicknesses for the front ETM (made of TiO2) and TCO (made of ITO) layers acting as a planar double-layer ARC [46]. The 500-nm perovskite layer in the planar reference exhibits high absorption along the visible spectrum, namely in the 375- to 750-nm wavelength range. Above such range there is an abrupt absorption drop when approaching the perovskite band gap (w1.55 eV). In the case of the thinner (250 nm) perovskite, there is a first smaller absorption drop at 600e750 nm wavelengths, and then above 750 nm there is a larger drop at the band gap. The incorporation of the LT structures improves the optical performance of the cells mainly by decreasing such drops in absorption observed for the planar PSCs, namely in the visible (600e750 nm wavelengths) for the thinner perovskite and in the NIR (>750 nm wavelengths) for both the 250 and 500 nm perovskites. Such absorption enhancement is due to a better light in-coupling toward the perovskite layer provided by the front photonic structures, via two mechanisms: (1) Antireflection, due to the geometrical index matching caused by the pyramidal-like cross-sectional shape of the TiO2 features (see top of Fig. 9.2), which is favored by the similarity between the real parts of the

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refractive indices (nw2.5 in wavelength range of interest) of the TiO2 and perovskite media. (2) Forward scattering, since the TiO2 features with high-scattering cross section redirect the light rays to multiple angles toward the perovskite, which increases the optical path length of the poorly absorbed near-band gap light and can strongly trap it by coupling with waveguided modes confined in the cell. This effect leads to the sharp absorption peaks in the NIR in the red curves of Fig. 9.5A and B, and to the “hot spot” regions in the generation profiles (insets of Fig. 9.5A and B), formed due to the interference between the light waves propagating vertically and horizontally along the cell structure (see bottom of Fig. 9.2). The colored regions in the spectra of Fig. 9.5A and B, above the curves of the absorption in the perovskite, correspond to parasitic absorption occurring in the other materials of the cell (optical losses that do not generate photocurrent). The parasitic absorption in the planar references occurs chiefly in two spectral regions: (1) in the UV (300e400 nm wavelengths) since the imaginary part of the refractive index (k) of ITO and TiO2 increases abruptly for wavelengths below 400 nm; (2) in the NIR (>800 nm wavelengths) due to the free carrier light absorption in the ITO, which leads to an increase of its k for the longer wavelengths, and absorption at the cell rear in the HTM (made of Spiro-OMeTAD). In the PSCs with the LT structure the parasitic absorption occurs for the same reasons, but for shorter wavelengths it is more pronounced along a broader range (300e500 nm) extending to the visible spectrum, due to absorption in the relatively large TiO2 features at the cells’ front, as seen in the generation profiles of Fig. 9.5A and B. In the NIR, the parasitic absorption takes place mainly in the top ITO and in the bottom HTM, as occurred in the planar cells. The severe parasitic absorption caused by the front ITO explains why the optimization algorithm converged to the minimum allowed thickness (50 nm, for a reasonable electric contact) for such TCO layer. Contrary to the planar cases, the antireflection effects caused by the LT structures are chiefly given by the geometrical index matching that they provide on the cell front. So, the TCO layer does not provide any beneficial optical contribution in these structures, and the thinner it is the better. However, a TCO layer with at least 50-nm thickness must be considered for electrical purposes, as it acts as the front n-electrode.

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9.3.2.1 Light trapping plus UV blocking effect of photonic structures The main mechanism responsible for UV light degradation in PSCs is attributed to the photocatalytic activity of TiO2 occurring in the interface between the ETM and the perovskite. Upon prolonged UV exposure, the photogenerated holes react with oxygen radicals adsorbed at surface vacancies, which become deep traps within the perovskite increasing carrier recombination [85]. Therefore, the use of UV blocking front coatings capable of preventing the penetration of UV photons toward the PSC has proven to be beneficial, yielding longer time of stabilized efficiency [86,87]. However, such external coatings also provide some undesirable degree of shading for the longer wavelength (visible-NIR) photons, thus decreasing the cells’ photocurrent. Quite advantageously, as seen from the absorption spectra of Fig. 9.5A and B, the wave-optical TiO2 structures optimized for the PSCs’ front naturally act as a UV-protective layer that, at the same time, enhances the absorption of visible and NIR photons in the perovskite, thus leading to overall JPH gains while contributing to the device robustness under sunlight exposure. Fig. 9.5D represents such gains, relative to the planar reference PSCs, for the two LT geometries explored in Section 9.3.1: TiO2 domes and voids, and for both perovskite thicknesses (250 and 500 nm). Since the enhancements attained with the optimized void arrays are clearly superior to those of the domes, in this section we focused the discussion on the bestperforming void geometry. To evaluate the LT plus UV-blocking performance of the optimized designs, we also represent in Fig. 9.5D the gains considering the photocurrent density (JPH_VIS-NIR) determined only in the visible and NIR range (400e1000 nm wavelengths), i.e., excluding the main UV range (300e400 nm). Such JPH_VIS-NIR enhancements can be regarded as the improvements that would be achieved considering PSCs with a UV-blocking filter for improved stability. Although the highest absolute JPH (31.3 mA/cm2dindicated in Table 9.1) is attained with the TiO2 voids on the 500-nm-thick perovskite, the optimized LT structure on the 250-nm absorber yields the highest JPH gain. This is expected from the Lambertian limits analyzed in Section 9.2, which show that the lower the absorber thickness, the higher can be the absorption enhancement provided by optimized LT. The maximum values of JPH (33.3 and 35.3 mA/cm2, respectively, for 250- and 500-nm-thick perovskites) and LT enhancement (31% and 25%, respectively) determined with the theoretical Lambertian surfaces are close but above those attained in this work (see Table 9.1 and Fig. 9.5D), which indicates that there is still

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room for further improvement of the optical schemes in PSCs. The main reason for the lower photocurrents attained with the wave-optical structures, relative to the ideal Lambertian limits, is attributed to the parasitic losses in the front TCO and TiO2 materials. The most impactful result is the JPH gain (27% including UV and 34.3% within only visible-NIR range) attained with the optimized TiO2 voids on the 250-nm-thick perovskite, since it demonstrates an improved device architecture that can allow (1) higher photocurrent (hence efficiency) relative to conventional planar PSCs with thicker (500 nm) perovskite layers; (2) a potential threefold improvement in the device flexibility due to the thinner absorber, since the flexural rigidity of the cell scales with the third power of its total thickness; (3) a twofold reduction in the costs associated to the perovskite material; and (4) potentially better PSC stability due to the role of the LT structure as a UV-blocking filter. In addition, such enhancement is considerably higher than the state-of-the-art enhancement (14.4% with 300-nm-thick perovskite absorber) achieved thus far using conventional texturing on the FTO/glass substrate [51,52].

9.4 Integration of photonic structures via soft lithography In a cost-driven technological field such as PV, integrating photonic nano-/microstructures in solar cells should be accomplished with low-cost and scalable techniques compatible with industrial manufacturing. Therefore, soft lithography methods are preferable relative to conventional “hard” methods, such as UV photolithography or E-beam lithography, which present strong limitations in terms of costs and patterning area.

9.4.1 Colloidal lithography microfabrication There have been several recent advances regarding the fabrication of dielectric-based photonic structures in PV devices via soft lithography techniques [32,49]. In some works, nanoimprint lithography (NIL) techniques are used to construct the structures [20,88,89], using a mold that is imprinted on a deformable material. However, the scalability of the NIL approach is still a challenge, since large contact areas between the mold and the imprinted nano-/microstructures can result in sticking issues, and pattern fidelity is poor over large areas as the polymer chains in the stamping materials tend to elastically relax. In addition, NIL can only be applied to “soft” materials that can be molded and then cured at moderated temperatures, like

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polymers or materials synthesized from solegel procedures. Nonetheless, in photonic applications as those discussed in the previous sections, high refractive index media are often required. Therefore, photonic-structured transparent semiconductors or dense metal oxides are preferable since they can interact more strongly with light, thus enabling higher scattering cross sections and antireflection action. In this respect, colloidal lithography (CLdsee Fig. 9.6) is a much more versatile soft lithography technique, since it can pattern any material

Figure 9.6 Sketch of colloidal lithography (CL) methods to engineer photonic microstructures: CL Method 1d(1) A hexagonal array of colloidal polystyrene (PS) spheres (diameters of 1e2 mm) is patterned on the cell front. (2) Reactive ion etching (RIE) uniformly etches the particles (but not the underneath layers), increasing their spacing. (3) The dielectric (photonic) material is deposited, filling the interparticle spaces. (4) The particles are removed leaving an array of semispheroidal voidlike features [29,31]. CL Method 2dThe steps are similar to CL Method 1, but here the spheres act as mask to structure an underneath layer of the dielectric material, thus creating domelike front features. The RIE (step 2) has to be carefully tuned, to selectively etch the dielectric material and not penetrate toward the cell layers [19,90]. Before liftoff of the particles in step 4, a thin contact layer (e.g., a TCO) can be deposited in the front (step 3) to assist in the conductivity of the front electrode. The bottom-right profile illustrates the absorbed power density, modeled with FDTD, along a cross section of the cell structure. FDTD, finite-difference time domain; PVD, physical vapor deposition; TCO, transparent conductive oxide

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[30,31,39]. Therefore, CL has been applied to engineer not only photonic wavelength-sized elements, as described below, but also other innovative nano-/microstructures for distinct purposes such as metallic micromesh electrodes [91]. Importantly, CL has been demonstrated to allow low-cost and nanoscale-precise patterning throughout indefinitely large areas, thus being one of the best scalable approaches for the cost-effective integration of photonic structures in PV devices. Moreover, CL employs only lowtemperature steps (20%), inexpensive ( 175 mm in solution-grown CH3NH3PbI3 single crystals, Science 347 (2015) 967e970, https://doi.org/10.1126/science.aaa5760. [64] W.S. Yang, B.-W. Park, E.H. Jung, N.J. Jeon, Y.C. Kim, D.U. Lee, S.S. Shin, J. Seo, E.K. Kim, J.H. Noh, S. Il Seok, Iodide management in formamidinium-lead-halidebased perovskite layers for efficient solar cells, Science 80 (356) (2017) 1376e1379, https://doi.org/10.1126/science.aan2301. [65] S. Yang, W. Fu, Z. Zhang, H. Chen, C.-Z. Li, Recent advances in perovskite solar cells: efficiency, stability and lead-free perovskite, J. Mater. Chem. A. 5 (2017) 11462e11482, https://doi.org/10.1039/c7ta00366h. [66] L. Zheng, Y. Ma, S. Chu, S. Wang, B. Qu, L. Xiao, Z. Chen, Q. Gong, Z. Wu, X. Hou, Improved light absorption and charge transport for perovskite solar cells with rough interfaces by sequential deposition, Nanoscale 6 (2014) 8171e8176, https://doi.org/10.1039/C4NR01141D. [67] A.R. Pascoe, S. Meyer, W. Huang, W. Li, I. Benesperi, N.W. Duffy, L. Spiccia, U. Bach, Y.-B. Cheng, Enhancing the optoelectronic performance of perovskite solar

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CHAPTER TEN

Organic and perovskite photovoltaics for indoor applications Harrison Ka Hin Lee, Jérémy Barbé, Wing Chung Tsoi SPECIFIC, College of Engineering, Swansea University, Swansea, United Kingdom

10.1 Introduction Photovoltaic (PV) cells are well known for their outdoor application as solar cells. However, PV cells can also generate power by harvesting indoor light. They can be integrated with consumer electronics to power lowpower consumption electronic devices. One of the most familiar examples is PV cellepowered pocket calculators which were commercialized as early as the late 1970s [1e3]. Amorphous silicon PV cells are used to power liquid crystal displays as they are power efficient and capable of operating in the low voltage range of 1.5e2 V. More recently, there has been a growing interest in using PV cells which harvest indoor light to power electronic devices for the Internet-of-Things (IoT). Such PV-powered devices are attractive for use in smart homes, smart offices, and smart buildings [4e6]. Although silicon PV cells have been dominating the market of indoor light harvesting for power generation (indeed, probably the only type of PV devices commercially used for indoor application) [7], research studies on other types of PV materials for indoor applications are growing. These alternative technologies include IIIeV inorganic PV [8], dye-sensitized PV (DSPV) [9e12], organic PV (OPV) [13e16], and, more recently, perovskite PV (PPV) [17e19]. One direction of research is to increase the maximum power generation per unit area (Pmax) under indoor light illumination. IIIeV PV cells have a higher power conversion efficiency (PCE) than silicon PV cells under one sun (AM1.5G) illumination (standard testing condition for solar cells). As shown in Fig. 10.1A, they are also more efficient under indoor lighting: the Pmax of silicon PV under indoor light (1000 lux) is around 46.5 mW/ cm2 while GaAs and GaInP PV cells have Pmax as high as 80.5 mW/cm2 Solar Cells and Light Management ISBN: 978-0-08-102762-2 https://doi.org/10.1016/B978-0-08-102762-2.00010-0

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Figure 10.1 (A) Pmax of PV cells under 100e1000 lux indoor light illumination (data from Ref. [20] for Si, GaAs, and GaInP; Ref. [16] for OPV; Ref. [12] for DSPV; Ref. [19]; and Ref. [21] for PPV), (B) Evolution of Pmax of OPV and PPV cells under 1000 lux illumination in recent years (details of the data are listed in later sections in Tables 10.5 and 10.6, respectively). Note that the light sources used in different works are often different. DSPV, dye-sensitized photovoltaic; OPV, organic photovoltaic; PPV, perovskite photovoltaic; PV, photovoltaic.

and 92.6 mW/cm2, respectively [20]. However, as is the case for their one sun applications, IIIeV PV cells are also more expensive than silicon PV cells for indoor applications. The other research direction is to further reduce the cost of the PV cells for indoor applications. As is the case for one sun applications, solution processable PVs such as DSPV, OPV, and PPV, which do not require high temperatures and have less strict processing conditions, could achieve this aim. The year of 2018 was a fruitful year for the indoor performance of these types of PV, with Pmax under 1000 lux illumination reaching up to 78.3 mW/cm2, 101.1 mW/cm2, and 111.9 mW/cm2 for OPV, DSPV, and PPV cells, respectively! Fig. 10.1A summarizes the performance of these PV cells under indoor light. This shows that they are not only potentially cheaper, but also can have significantly better performances than silicon PV cells under indoor conditions. Fig. 10.1B (details are shown in later sections in Tables 10.5 and 10.6) shows the reported indoor performance of OPV and PPV under 1000 lux illumination in recent years. The aim of this chapter is to introduce the emerging field of using OPV and PPV cells to harvest indoor light for power generation. Both types of device have shown promising potential in this direction. Besides being solution-processable, both their optical (e.g., absorbance) and electronic properties (frontier orbitals/electronic bands) can be tuned by chemical

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approaches. This allows materials to be easily adapted to better suit indoor conditions and hence enhance indoor performances. Furthermore, flexible and light weight devices can be fabricated on flexible plastic substrates. Interestingly, although OPV and PPV cells are less stable than silicon PV cells under one sun conditions, their lifetime under indoor lighting could be much longer due to the significantly less harsh conditions (e.g. light level, temperature and humidity) in an indoor environment compared to an outdoor environment. Moreover, as the series resistance is much less crucial for low light applications, owing to the much lower current flowing through the devices compared to outdoor operation, the layers in the devices could be made thicker, which is particularly beneficial for scaling up the fabrication of OPV devices. This chapter is structured as follows: Section 10.2 provides basic principles on PV cells for indoor applications, particularly the optimum band gap (vs. that for outdoor applications), Section 10.3 specifies characterization methods relevant to indoor performance, particularly the light sources used for indoor lighting. Sections 10.4 and 10.5 introduce the performance of OPV and PPV cells under indoor light, respectively. Section 10.6 gives a brief introduction to PV cells for indoor applications, particularly about their use within IoT devices, which are rapidly growing in number and are key elements for smart homes, smart offices, and smart buildings. Finally, a summary and outlook of OPV and PPV cells for indoor applications is provided in Section 10.7.

10.2 Basic principles for indoor photovoltaics One main difference between AM1.5G and indoor light is the spectrum width. AM1.5G solar light has a much broader range (300e2500 nm) including ultraviolet, visible, and infrared light (see Fig. 10.2A). Energy efficient indoor lighting emits photons mainly within the visible range (ca. 400 nme700 nm), which considers the response of human eyes. The most common lighting technologies nowadays are fluorescent lamps (a typical spectrum is shown in Fig. 10.2A) and light-emitting diodes (LEDs). Both types of lighting are much more energy efficient than incandescent lighting where the emission follows the blackbody radiation curve, covering not only the visible region but also the infrared region. The infrared region does not make any contribution to the illumination, due to the insensitivity of our eyes to infrared light. Rather, it primarily heats up the surroundings.

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Figure 10.2 (A) Spectra of AM1.5G illumination and a typical fluorescent lamp. (B) Simulated maximum PCE versus Eg of PV materials under different light sources. PCE, power conversion efficiency; PV, photovoltaic. Reproduced from M.F. M€ uller, M. Freunek, and L.M. Reindl, “Maximum efficiencies of indoor photovoltaic devices,” IEEE J. Photovoltaics, 3 (1) (2013). 59e64. with permission from IEEE.

The theoretical limit of the PCE of PV cells under one sun illumination can be calculated based on the energy gap (Eg) of the PV materials. Fig. 10.2B shows a simulation plot of PCE against Eg of PV materials under different light sources, including AM1.5G, fluorescent lamps, and LED light sources [22]. For AM1.5G illumination, it is well known that the limit of PCE is 33% where the PV materials have an Eg of 1.1e1.4 eV, corresponding to an optical absorption cut off of 1127 e886 nm, respectively. On the other hand, PV materials with an Eg of 1.8e1.9 eV (corresponding to an optical absorption cut-off at ca. 680 nm) can reach a PCE limit of 50%e60% under indoor lighting. The Eg of PV materials for indoor light harvesting is larger as only photons from the visible region need to be collected. The maximum voltage that can be generated in a PV cell is called opencircuit voltage (VOC), which can be expressed as [23].   Iph nkT VOC w (10.1) ln e I0 where n is diode ideality factor, k is the Boltzmann constant, T is the absolute temperature, e is the elementary charge, Iph is the photocurrent, and I0 is the reverse saturation current of the diode. From Eq. (10.1), the VOC depends on several parameters such as the diode ideality factor, which describes how ideal the diode is (n ¼ 1 for ideal diode), light intensity which linearly affects Iph, and temperature. For a typical indoor environment, the light intensity is about two to three orders of magnitude lower (0.1e1 mW/cm2) than the

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one sun illumination (100 mW/cm2). As a result, the VOC obtained in an indoor environment is always lower than that under one sun illumination. Table 10.1 shows the simulated percentage drops in VOC (with different VOC values under one sun) when the light intensity is reduced to 0.1 mW/ cm2 (1000 times lower than one sun), for ideality factors of 1, 1.5, and 2. Since the PCE is equal to the product of JSC, VOC, and fill factor (FF) divided by the incident light intensity (Pin), the drop in VOC directly affects the performance. To reduce the percentage drop, a PV system with a high VOC at one sun and a low ideality factor is more favorable; such a system is therefore preferred for low light applications. The FF is a complicated parameter which can be affected by several factors such as charge generation and recombination, shunt resistance (RSh), series resistance (RS), etc. [24]. Although the charge generation and recombination are light dependent, different PV systems may behave differently at low or high light levels. For OPV, Steim et al. simulated the JeV characteristic of a classical polymer:fullerene bulk-heterojunction system, P3HT:PCBM, at one sun and low light conditions with different parasitic RS and RSh [14]. They found that a high RS limits the performance at high light levels, while a low RSh limits the performance at low light levels. Shunt resistance is closely related to leakage current [25]; a low RSh results in a high leakage current and vice versa. Since Iph is much lower under low light conditions, the magnitude of the leakage current can be easily comparable or even higher than Iph. As a result, the overall current will be affected or dominated by the leakage current. Moreover, when the leakage current is high (a low RSh), Eq. (10.1) is no longer valid as it is based on the assumption that RSh is high. In this case, the VOC drops even more at low light intensities due to the high leakage current [25]. Lechêne et al. systematically studied how the VOC of an OPV cell drops versus decreasing light intensity with different RSh values by adjusting the thickness of the electron transporting layer (ETL) [26]. All these findings suggest that low leakage current (or high RSh) is essential for PV cells to achieve high PCEs at low light levels.

10.3 Characterization of indoor photovoltaic cells 10.3.1 Indoor light sources For energy-efficient indoor lighting, fluorescent lamps and LEDs which mainly emit visible light are used (see Fig. 10.3 for a photo of a fluorescent lamp array and its emission spectrum). Illuminance is a measure of

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Table 10.1 VOC at 100 mW/cm2 and 0.1 mW/cm2, and the associated percentage drops in VOC with an ideality factor of 1, 1.5, and 2, respectively. VOC (V) at VOC (V) at 0.1 mW/ Percentage drop VOC (V) at 0.1 mW/ Percentage drop VOC (V) at 0.1 mW/ Percentage drop 100 mW/cm2 cm2 if n [ 1 cm2 if n [ 1.5 cm2 if n [ 2 in VOC (%) in VOC (%) in VOC (%)

1.4 1.2 1.0 0.8 0.6 0.4

1.22 1.02 0.82 0.62 0.42 0.22

12.7 14.8 17.7 22.2 29.6 44.4

1.13 0.93 0.73 0.53 0.33 0.13

19.0 22.2 26.6 33.3 44.4 66.5

1.05 0.85 0.65 0.45 0.25 0.05

25.4 29.6 35.5 44.4 59.2 88.7 Harrison Ka Hin Lee et al.

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Figure 10.3 (A) The photo shows an array of fluorescent lamps illuminating an organic photovoltaic module (a lux meter is also shown), (B) a spectrum of the fluorescent lamps at 300 lux.

light level in which the eye sensitivity as a function of wavelength is taken into consideration. The unit for illuminance is lux. Table 10.2 shows the typical illuminance in various environments. For most indoor environments, the illuminance ranges from 100 lux to 1000 lux. According to the European Standard EN 12464-1:2002 (Light and lighting - Lighting of work places - Part 1: Indoor work places), for a normal office environment, 300e500 lux is suggested. However, unlike the standard one sun condition, to date there is no common standard for testing indoor PV cells. Nevertheless, performance between 200 lux (a dim indoor environment) and 1000 lux (a bright indoor environment) is usually reported.

10.3.2 Reporting indoor PV performance To study the indoor PV performance, Pmax is usually reported at a given lux level which can be extracted from the corresponding JeV curve obtained under such illumination. Pmax is the product of the voltage and current

Table 10.2 Typical illuminance in various environments. Illumination condition

Illuminance (lux)

Full moon Street lighting Home lighting Office desk lighting Surgery lighting Direct sunlight

1 10 30e300 100e1000 10000 100000

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density at the maximum power point (Vmax, Jmax), which can also be calculated by JSC  VOC  FF where FF ¼ (Jmax  Vmax)/(JSC  VOC). For a fair comparison between reported Pmax values, the use of an accurate lux meter is essential. It is important to note that some lux meters may be crudely made, which could lead to inaccurate measurements. On the other hand, there are some well-calibrated lux meters by third parties, such as National Institute of Standards and Technology, which are preferable for more reliable measurements. In addition to the Pmax, the PCE is sometimes reported in the literature. PCE provides a more precise indication on how efficiently the incident light is converted into output power. To obtain the PCE, a measurement of the incident light intensity (Pin) is required, as PCE ¼ Pmax/Pin. However, an accurate measurement of Pin is challenging as the light intensity of the indoor sources is typically quite low, about three orders of magnitude less than the AM1.5G illumination, and thus a sensitive spectroradiometer or power meter is required for accurate evaluation. Furthermore, for PPV and DSPV, it is quite common to observe hysteresis in the JeV scans, i.e., the shape of the JeV curves depends on the scan direction (forward or reverse). Additionally, the scan rate may also affect the shape of the JeV curves. Chen et al. conducted an intense study on the variation of indoor performance of DSPV cells involving 15 research groups around the world [28]. Their findings show that the hysteresis could be more serious under low light conditions and suggested that a low scan rate is preferable, and that stabilized power generation measurements (Pmax over time) would be a more accurate way to probe Pmax.

10.4 Organic photovoltaic cells for indoor light harvesting 10.4.1 Background of organic photovoltaic cells PV characteristic of organic materials was first observed in the 1950s [33]. With continuous efforts over the past few decades, PCEs have now reached over 15% for single junction devices [34] and over 17% for tandem devices [35]. Although OPV cells are unlikely to be able to achieve comparable PCEs to inorganic PV cells under one sun illumination, OPV cells have some unique features which other PV technologies may not have. Some benefits include low device fabrication costs, compatibility of fabrication with flexible substrates, high-specific powers, and highly tuneable

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material properties, which, among other things, allow the fabrication of semitransparent and visibly transparent devices [33]. The active layer within OPV devices usually consists of two organic materials, one electron donor and one electron acceptor, forming a bulkheterojunction blend layer. Over the past 20 years, the main developments in OPV have been driven by new donor materials. For example, P3HTbased devices demonstrated 2%e5% PCE, PCDTBT-based devices demonstrated 5%e7% PCE, PTB7 or PTB7-Th-based devices demonstrated 7% e10% PCE, and PffBT4T-2OD (PCE11)-based devices demonstrated 11% PCE. These donor materials perform well when they are blended with fullerene derivatives such as PCBM or PC71BM, which have been the main acceptor materials for OPV over the past two decades. Recently, there has been a breakthrough in the development of new acceptor materials which are classified as small molecule nonfullerene acceptors (NFAs). NFAs allow much more tunability in optical and electrical properties of the acceptor materials leading to a significant increase in PCEs from ca. 11% to 15%.

10.4.2 Indoor performance of three benchmark OPV systems OPV cells tested under indoor light environments were reported in 2011 [14]. At that time, the focus was on the P3HT:PCBM system as it had a good PCE under one sun illumination (PCE of 2%e5%) [36]. However, the indoor performance of this system is not that attractive with PCE of 7% (Pmax ¼ 19 mW/cm2) under 1000 lux illumination [14]. Although more polymer:fullerene systems with higher PCEs have been developed later, there are only limited works revisiting the applicability of OPV cells for indoor applications. In this section, the AM1.5G and indoor performance of three benchmark polymer:fullerene systems (P3HT:PCBM, PCDTBT:PC71BM, and PTB7:PC71BM) are compared as a case study. Through this comparison, we can examine the criteria for selecting OPV systems for indoor applications (see Section 10.2). The devices were first tested under AM1.5G illumination as a reference and then tested in an enclosed light box which consists of an array of commercially available fluorescent lamps with a reflector surrounding them to provide a uniform lighting environment. Fig. 10.4 shows the JeV curves of the three types of OPV cell under solar simulator and fluorescent lamp illumination at 300 lux. The device parameters under these illumination conditions are shown in Table 10.3.

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Figure 10.4 JeV characterization of P3HT:PCBM (black square), PCDTBT:PC71BM (red circle), and PTB7:PC71BM (blue triangle) OPV devices under (A) AM1.5G conditions with an intensity of 82e90 mW/cm2 and (B) 300 lux fluorescent lamp illumination. PCE, power conversion efficiency; OPV, organic photovoltaic. Reproduced from H.K.H. Lee, Z.Li, J.R. Durrant, W.C. Tsoi, Is organic photovoltaics promising for indoor applications?, Appl. Phys. Lett., 108(25) (2016) with permission from AIP.

The PTB7:PC71BM device shows the highest PCE of 6.8% under AM1.5G illumination followed by the PCDTBT:PC71BM and P3HT:PCBM devices with PCEs of 6.0% and 2.4%, respectively. To calculate the PCE under the fluorescent lamps, the Pin at 300 lux was measured as 83.6 mW/cm2. The Pmax values of these three systems under 300 lux are 4.8 mW/cm2, 13.9 mW/cm2, and 12.2 mW/cm2 for P3HT:PCBM, PCDTBT:PC71BM, and PTB7:PC71BM, which correspond to PCEs of 5.8%, 16.6%, and 14.6%, respectively. Importantly, the PCEs of all the systems under 300 lux fluorescent lamp illumination are substantially higher than their PCEs under the solar simulator since a larger fraction of the light can be converted into electricity. A large portion of photons (ca. 40% of the photons) from the infrared region of the solar spectrum are not absorbed in these devices while most of the photons from the fluorescent lamps, which emit mainly visible

Table 10.3 Device parameters of P3HT:PCBM, PCDTBT:PC71BM, and PTB7:PC71BM OPV cells under 300 lux fluorescent lamp illumination. JSC (mA/cm2) VOC (V) FF (%) Pmax (mW/cm2) PCE (%)

P3HT PCDTBT PTB7

20.6 27.7 28.6

0.41 0.72 0.61

56.6 69.3 69.5

4.8 13.9 12.2

5.8 16.6 14.6

Reproduced from H.K.H. Lee, Z.Li, J.R. Durrant, W.C. Tsoi, Is organic photovoltaics promising for indoor applications?, Appl. Phys. Lett., 108(25) (2016) with permission from AIP.

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light, are absorbed. The PTB7:PC71BM device has a higher JSC than the PCDTBT:PC71BM device under AM1.5G illumination as the former has a broader external quantum efficiency (EQE) spectrum covering up to 750 nm, while the EQE of the latter only covers up to 700 nm [37,38]. On the contrary, both types of devices have similar JSC under fluorescent lamps, as the number of photons with a wavelength over 700 nm is very low. Therefore, due to the higher VOC of the PCDTBT:PC71BM devices, the PCDTBT:PC71BM cells have higher PCE under the fluorescent lamp than the PTB7:PC71BM cells. The Eg of P3HT, PCDTBT, and PTB7 are 2.0 eV, 1.9, and 1.75 eV, respectively [36e38], with the Eg of PCDTBT matching the best with the optimal Eg for harvesting indoor light (1.8e1.9 eV). A PV material with this Eg can absorb most of the visible light and at the same time maintain a high VOC. This correlates well with the highest VOC obtained for the PCDTBT:PC71BM device (among the investigated devices) under both the AM1.5G and indoor lighting conditions. Fig. 10.5 shows the VOC as a function of light intensity for all three devices. By using Eq. (10.1), the diode ideality factor, n, of the devices can be extracted from the gradient of these plots. The extracted ideality factors of the devices are 1.25, 1.24, and 1.10 for the P3HT:PCBM, PCDTBT:PC71BM, and PTB7:PC71BM devices, respectively. This is consistent with the literature where most optimized OPV systems have n close to 1. A similar n suggests that the magnitude of the drop in VOC should be similar when light intensity is reduced across a

Figure 10.5 VOC (solid symbol) and fill factor (FF) (open symbol) against light intensity under AM1.5G illumination. The two data points of the PCDTBT:PC71BM device at the lowest intensity are not included in the fit. Reproduced from H.K.H. Lee, Z.Li, J.R. Durrant, W.C. Tsoi, Is organic photovoltaics promising for indoor applications?, Appl. Phys. Lett., 108(25) (2016) with permission from AIP.

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given range. Since the PCDTBT:PC71BM device has the highest VOC among these systems, the percentage drop in the VOC from high light to low light intensity is the least. For instance, when the light intensity was reduced from w90 mW/cm2 to 0.1 mW/cm2, the percentage drop in VOC of PCDTBT:PC71BM and P3HT:PCBM devices is around 23% and 38%, respectively. This result correlates well with the predicted trend in VOC shown in Table 10.1. Note that the two lowest VOC data points for PCDTBT:PC71BM devices are slightly off the fitting line. This phenomenon is often observed in devices when the leakage current is not low enough. Nevertheless, the leakage current in these devices should not be too serious even for the PCDTBT:PC71BM devices. Lechêne et al. found that the thickness of the ETL (namely PEIE) used in an inverted device structure could have strong influence on the leakage current. Thin PEIE layer causes a high leakage current and thus significantly affects the VOC at low light intensities. Based on their findings, they suggested that the leakage current should be less than w10 mA/cm2 if the incident light intensity is around 0.1 mW/cm2. Fig. 10.5 also shows that generally the FFs of the three different devices increase with decreasing light intensity. This can be attributed to lower charge recombination and the smaller effect of RS at lower light levels [24]. In particular, the FFs of PCDTBT:PC71BM and PTB7:PC71BM devices can reach around 70% in the intensity range of 0.1e1 mW/cm2 (approximately the indoor light intensity). Therefore, OPV devices could be more favorable for low light operation in terms of FF. In summary, this study of the indoor performance of the three benchmark OPV systems verifies the following criteria for indoor PV applications: (1) Eg of the PV materials should be 1.8e1.9 eV. (2) A high VOC system (at one sun) is more beneficial provided that n is close to 1. (3) Leakage current of the device should be sufficiently low. As demonstrated, the PCDTBT:PC71BM cell performs better than the PTB7:PC71BM cell under the low light conditions, which suggests that the highest performing system under AM1.5G illumination is not necessarily the best under indoor lighting conditions.

10.4.3 Finding state-of-the art OPV systems for indoor applications After the generation of PCDTBT:PC71BM and PTB7:PC71BM, there were a number of new donor materials synthesized with higher one sun PCEs

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[39,40]. The strategy to enhance the PCE is to reduce the Eg of the donor materials to achieve higher JSC. Moreover, recently, several successful NFAs have been synthesized which have allowed record PCEs for OPV to be reached. However, most of these donors and NFAs have an Eg smaller than 1.8 eV which is not ideal for the indoor applications. Nevertheless, we have found that a system, first reported in 2015, consisting of a small molecule donor, namely BTR, and fullerene acceptor, PC71BM, which fulfills all the criteria for indoor applications [41]. BTR:PC71BM is a solution processable system with an optimized active layer thickness of ca. 250 nm with a high FF of over 75% and a PCE of 9.3%. Furthermore, the system can maintain 90% of the optimal PCE even with a 400-nm-thick active layer. More importantly, BTR has an Eg of 1.8 eV, and the BTR:PC71BM device has VOC > 0.9 V with high EQE in the visible light region. This suggests that BTR:PC71BM should be an excellent system for indoor light harvesting. Here, we explore this BTR:PC71BM system through optimization on solvent vapor annealing (SVA) for both one sun and indoor conditions. The JeV curves of the BTR:PC71BM devices obtained under AM1.5G illumination are shown in Fig. 10.6A for different SVA treatment times. The device parameters obtained under AM1.5G illumination and indoor lighting are summarized in Table 10.4. Both the JSC and FF show maxima at an SVA treatment time of 2e5 min. A gradual decrease in the VOC from w1 V to a saturated value of w0.92 V was observed with increasing SVA time, and the optimized PCE was >10% under AM1.5G illumination. All these devices were then tested under the fluorescent lamps at 200 and 1000 lux, which corresponds to a Pin of 55.7 mW/cm2 and 278.7 mW/cm2, respectively. Fig. 10.6B shows the JeV curves of the devices under the indoor illumination with the representative treatment times. The device with 2 min treatment shows the highest Pmax of 14.6 mW/cm2 at 200 lux (PCE ¼ 26.2%) and a Pmax of 78.2 mW/cm2 at 1000 lux (PCE ¼ 28.1%). The EQE spectra shown in Fig. 10.6C demonstrate excellent conversion across the visible spectrum. The Pmax at different illuminance is plotted in Fig. 10.6D for the BTR:PC71BM device with the optimized SVA time of 2 min. In addition to the works mentioned above, there are further published works using OPV cells for indoor PV. To have an overall picture of the development of the OPV for indoor applications in recent years, we have summarized the representative indoor performance of OPV devices from the literature in Table 10.5 for reference.

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Figure 10.6 JeV curves of BTR:PC71BM devices at various SVA times, measured under (A) AM1.5G solar simulator and (B) fluorescent lamp illumination. (C) EQE spectra of the devices with 0, 2, and 30 min of SVA treatment. The spectrum of the fluorescent lamps is also shown as a reference (in arbitrary scale). (D) Pmax of the device with 2 min SVA at different illuminance. EQE, external quantum efficiency; SVA, solvent vapor annealing. Reproduced from H.K.H. Lee et al., Organic photovoltaic cells e promising indoor light harvesters for self-sustainable electronics, J. Mater. Chem., 6, (14) (2018) 5618e5626. With permission from the Royal Society of Chemistry.

10.5 Perovskite photovoltaic cells for indoor light harvesting 10.5.1 Perovskite semiconductors PPV cells have been one of the hottest topics in solar cells in recent years because of the remarkable optical and electronic properties of the perovskite semiconductors and hence rapid progress in PCE. In just a few years, their maximum PCE increased from 3.8% in 2009 [42] to certified 23.7% [43], which was achieved after optimization of the device architecture (including standard n-i-p and inverted p-i-n architectures), perovskite composition, interlayers, and fabrication processes.

0 0.5 1 2 5 10 30

11.6 [118.1] 11.8 [117.8] 13.0 [127.0] 13.3 [133.1] 12.5 [121.8] 11.5 [108.8] 9.7 [98.1]

0.994 0.995 0.966 0.945 0.928 0.923 0.922

[0.856] [0.859] [0.815] [0.791] [0.755] [0.747] [0.754]

51.6 53.5 68.9 75.1 79.1 75.4 61.9

[67.4] [66.3] [73.6] [75.2] [77.7] [75.6] [73.7]

5.97 6.26 8.65 9.42 9.19 7.98 5.54

[68.0] [66.3] [74.8] [78.2] [71.0] [60.9] [53.5]

6.6 [24.2] 7.0 [23.8] 9.6 [27.0] 10.5 [28.1] 10.2 [25.3] 8.9 [21.8] 6.2 [19.3]

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Table 10.4 Device parameters for different SVA times, measured under AM1.5G with 90 mW/cm2 intensity. Values in the brackets are data obtained under 1000 lux fluorescent lamps. Pmax (mW/cm2) [mW/ 2 2 cm2] VOC (V) FF (%) PCE (%) SVA time (min) JSC (mA/cm ) [mA/cm ]

FF, fill factor; PCE, power conversion efficiency; SVA, solvent vapor annealing. Reproduced from H.K.H. Lee et al., Organic photovoltaic cells e promising indoor light harvesters for self-sustainable electronics, J. Mater. Chem., 6, (14) (2018) 5618e5626. With permission from the Royal Society of Chemistry.

369

2011 P3HT:PCBM (C, 1) 2014 a-6T/Subnc/SubPc (C, 0.134) 2015 PTB7-Th:PC71BM (C, 1)

FL LED

1000 N/A

270 N/A

19 N/A

7 20.45

19 N/A

[14] [14a]

LED

10.55 11.63 11.5 8.7 0.6 5.8 16.6 14.6 11.2

69.0b 30b 12.7b 16.0b 46.3b 40.7c 31.3c

[26]

300

7.6a 42.3a 20.7 9.0 3.8 4.8 13.9 12.2 9.38

[14b]

FL LED ILB FL

72.3 364 180 104 633 83.6

47.6b

2016 PCDTBT:PC71BM (C, 0.475)

186 890 300

LED

N/A

N/A

N/A

12.83c 21.04c

N/A

[15a]

FL LED LED

1000

N/A

N/A

186.6a

6.7

44.37 54.76 25.0b

[15b]

500

44.37 54.76 12.5

FL LED FL LED FL LED

500

164a 176a 164a 176a 164a 176a

15.77c 15.67c 22.57c 22.97c 21.56c 23.23c

9.59c 8.90c 13.76c 13.05c 13.14c 13.20c

32.59c 32.13c 48.09c 48.42c 41.90c 44.09c

[63a]

2016 P3HT:PCBM PCDTBT:PC71BM PTB7:PC71BM (C, 0.15) PCDTBT:PC71BM (M, 100) 2016 P3HT:PCBM PTB7-Th:PC71BM (C, N/A) 2017 N/A (M, 3.6664)

P3HT:ICBA PTB7-Th:PC71BM (C, N/A)

[15]

[64]

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2017 PBTZT-stat-BDTT-8: PCBM:PC71BM (24% transmission) (C, 15.5) 2017 P3HT:PCBM

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Table 10.5 Indoor performance of organic photovoltaic devices from representative works. Pmax Pmax at 1000 System (C/M, Light source Illuminance (lux) Pin (mW/cm2) (mW/cm2) PCE (%) lux (mW/cm2) References Year area in cm2)

FL

2018 PBDTTT:PC71BM (C, N/A) 2018 PCDTBT:PDTSTPD: PC71BM (C, N/A) 2018 P3HT:ICBA (C, 0.1) 2018 P3HT:ICBA (C, 0.1) 2018 P1:PC71BM PCDTBT:PC71BM (C, N/A) 2019 BDT-2T-ID:PNP BDT-1T-ID:PNP PTB7:PNP (C, 0.06) BDT-2T-ID:PNP (M, 9.5)

FL LED FL LED LED LED LED

LED

200 500 1000 1000

14.6 38.4 78.2 53c 55c 15.4 14.7 22.7c 30.7a,c 14.86 14.53

26.2 27.6 28.1 16.16c 15.58c 20.8 18.9 13.4c 5.4c 19.15 18.72

78.2

[16]

53c 55c 51.3b 49.0b 45.4b,c 30.7a,c 49.5 48.4

[16a]

500 1000 300

55.7 139.3 278.7 328a 353a 74.0 77.6 170 568a 77.6

200

76.8

12.3c 9.5c 7.3c 11.7a

16.0c 12.4c 9.5c 15

61.5b,c 47.5b,c 36.5b,c 58.4

300

C, cell; FL, fluorescent lamp; ILB, incandescent light bulb; LED, light-emitting diode; M, module; PCE, power conversion efficiency. a Extracted from the other reported values. b Linearly scaled from the reported data. c Averaged value.

[16b] [16c] [16d] [16e]

[16f]

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The term “perovskite” refers to the mineral CaTiO3 and any structure adopting the same ABC3 structural framework [44]. In the case of perovskite semiconductors, A is an organic or inorganic cation (MA: CH3NHþ 3 , FA: þ þ CHN2Hþ , Cs , or Rb ), B is a metallic cation, and C is a halogen (Cl, 4 Br or I). Depending on their composition, perovskites exhibit a wide range of optical and electronic properties, which can be optimized for different applications. There is a large class of perovskite semiconductors with promising PV characteristics. The most investigated is methylammonium lead iodide CH3NH3PbI3 (or MAPbI3). However, in recent years, mixed ion perovskites (FAxMA1-xPbBryI3-y) have attracted more and more attention due to their superior PCE and stability [45]. Another chapter in this book provides more details on the properties and fabrication of perovskite solar cells [46]. Beside high absorption coefficients of 105 cm1 [47] and long electrone hole diffusion lengths [48,49], one of the main advantages of perovskite semiconductors for indoor PV applications is the ability to tune their band gap Eg by ion substitution. This tunability enables improved matching with the indoor light spectrum, and maximization of the current and voltage generated, i.e., the output power. As an example, iodide in CH3NH3PbI3 can be replaced by halides with smaller ionic radii (chloride or bromide) to increase the Eg. For single crystal perovskite semiconductors, Eg values of 2.97, 2.24, and 1.53 eV have been measured for MAPbCl3, MAPbBr3, and MAPbI3, respectively [50]. Intermediate Eg values could be achieved simply by mixing the halides in various ratios. For example, an optimum Eg of 1.8e1.9 eV for indoor lighting could be obtained by mixing iodide and bromide in the MAPbI3xBrx film with a ratio of 1:1 [51].

10.5.2 Architectures of PPV cells Traditionally, PPV cells have been designed with the use of a mesoporous scaffold as shown in Fig. 10.7A. Standard mesoporous PPV (mPPV) cells are commonly composed of a transparent conductive oxide, a compact titanium dioxide (TiO2) hole blocking layer, a mesoporous TiO2 ETL, a perovskite PV material infiltrated into the mesoporous scaffold, an HTL (the most commonly used is Spiro-OMeTAD, and a metal electrode. Standard planar PPV cells (Fig. 10.7B) use a similar architecture but without the mesoporous scaffold. By using a compact tin oxide (SnO2) ETL instead of TiO2, a high PCE can be attained with the standard planar PPV structure, almost as high as the mesoporous counterpart [52]. Another planar configuration has been successfully developed where the layer stack is inverted (TCO/HTL/

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Figure 10.7 Schematic diagram of four types of perovskite photovoltaic (PPV) device architectures. (A) Standard mesoporous PPV; (B) standard planar PPV; (C) inverted planar PPV; (D)carbon-based mesoporous PPV.

Perovskite/ETL/Metal) (iPPV), as shown in Fig. 10.7C. In this case, PEDOT:PSS and PCBM are commonly used as the HTL and ETL, respectively, but a wide range of alternative interlayers have been developed which have the potential to increase the PCE and improve the device stability. Finally, a carbon-based mesoporous and HTM-free PPV (cPPV) configuration has been developed (Fig. 10.7D). The cell architecture is based on a stack of mesoporous TiO2 ETL, a mesoporous zirconium dioxide (mZrO2) spacing layer, and a thick mesoporous carbon electrode, in which the perovskite solution is infiltrated. In this case, as the entire stack is printable, relatively stable, and does not include a metal top electrode, it is believed to be the readiest system for large-scale production. However, it still suffers from a lower PCE due to the high series resistance of the carbon electrode.

10.5.3 Introduction to perovskite PV cells for indoor applications The first study investigating the indoor PCE of PPV was reported in 2015 using inverted cells with the structure PEDOT:PSS/MAPbI3/PCBM or C60/TmPyPB/Ag [17]. After optimizing the ETL deposition process, the best-performing small area device exhibited a Pmax of 83.4 mW/cm2 with no hysteresis behavior, under fluorescent lamp illumination at 1000 lux. More recently, a Pmax of 111.9 mW/cm2 was achieved at 1000 lux for iPPV using Poly-TPD as an HTL instead of the more commonly used PEDOT:PSS [21]. At 200 lux, the Pmax maintained a relatively high value of 19.6 mW/cm2 for these cells. This study shows the importance of choosing the right ETL/HTL materials for achieving high performance under indoor lighting.

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For PPV devices with the standard structure m-TiO2/MAPbI3/SpiroOMeTAD/Au, an impressive Pmax of 115.6 mW/cm2 was measured under 1000 lux of fluorescent lamp illumination, with 19.9 mW/cm2 at 200 lux [21]. In another study, a slightly higher Pmax of 20.1 mW/cm2 was obtained at 200 lux of a white LED illumination using a SnO2/MgO composite ETL instead of m-TiO2 [19]. The indoor performance of cPPV was also reported. Despite these cells having a significantly lower PCE under one sun, mostly because of the high series resistance of the thick carbon electrode, they exhibited high Pmax of 26.6 and 89.4 mW/cm2 under 200 and 1000 lux fluorescent lamps, respectively. Although the high RS value leads to a low FF (58%) under one sun, it does not cause a low FF under low light (FF ¼ 71.9% for 1000 lux and 70.4% for 200 lux). Hence it was suggested that devices having high RS, like the cPPV, perform better at low light levels. Due to these promising results, together with its low cost and relatively good stability, it was envisioned that this device architecture might be the closest to commercialization for indoor applications. Although most of these studies were carried out on small area devices (usually 0.1 cm2), Chen et al. also fabricated a larger PPV device with an area of 5.44 cm2 based on the inverted structure ITO/PEDOT:PSS/ MAPbI3/PCBM or C60/TmPyPB/Ag [17]. This device exhibited a steady power of 340 mW under 1000 lux of fluorescent lamp illumination which is considered readily sufficient for modern sensing nodes of the IoT. In another study, a cPPV module with active surface area of 9.75 cm2 was fabricated and exhibited a high power output of 683 mW at 1000 lux of fluorescent lamp illumination (Pmax ¼ 70.1 mW/cm2), which is significantly better than silicon PV modules (Pmax of w15e46 mW/cm2 under 1000 lux) [20,53,54]. An all-printable large A4-sized module was also fabricated using the carbon-based architecture, which yielded z2 mW at 200 lux, 6 mW at 600 lux, and 10 mW at 1000 lux of fluorescent lamp illumination [55]. In general, the tolerance of a high RS under low light could imply more flexibility on the design of the interconnections between subcells in a module [21]. Table 10.6 summarizes the indoor performance of PPV devices from representative works.

10.5.4 Perovskite PV cell operation in low-light environments Despite these initial results showing the remarkable performance of PPV for indoor light harvesting as compared to commercially available devices based

2015 (Inverted1, 0.051) (Inverted1, 0.051) (Inverted1, 0.051) (Inverted1, 5.44) 2016 (Normal1, 0.12) (Normal1, 0.12) 2017 (Flexible1, 0.1) (Flexible1, 0.1) 2018 (Normal2, 0.1) (Normal2, 0.1) 2018 (Carbon1, 198) (Carbon1, 198) (Carbon1, 198) (Carbon1, 198) 2018 (Normal3, 0.09) (Normal3, 0.09) (Inverted2, 0.045) (Inverted2, 0.045) (Carbon1, 0.49) (Carbon1, 0.49) (Carbon1, 9,75)

FL

FL LED LED FL

FL

100 600 1000 1000 200 400 200 400 200 400 100 200 600 1000 200 1000 200 1000 200 1000 1000

31.7a 194.7a 316.8a 316.8a 64.2a 128.3a 76.0a 144.1 80.8a 154.6a 28a 57a 169a 280a N/A N/A N/A N/A N/A N/A N/A

6.6a,c 48.9a,c 83.3a,c 62.5a,c 15.4 32.6 9.77 19.2 20.2 41.6 3.03a 9.09a 28.8a 50.5a 19.9 115.6 19.6 111.9 16.3 89.4 70.1

20.9c 25.1c 26.3c 20.4c 24.0 25.4 12.85 13.32 25.0 26.9 11 16 17 18 N/A N/A N/A N/A N/A N/A N/A

References

[17] 83.3a,c 62.5a,c 81.5b

[18]

48b

[18a]

104c

[19]

50.5a

[55]

[21]

Organic and perovskite photovoltaics for indoor applications

Table 10.6 Indoor performance of perovskite photovoltaic devices from representative works. Illuminance Pmax at 1000 lux (mW/ Perovskite cell Pin (mW/cm2) Pmax (mW/cm2) PCE (%) cm2) Year (structure, area in cm2) Light source (lux)

115.6 111.9 89.4 70.1

375

FL, fluorescent lamp; LED, light-emitting diode; PCE, power conversion efficiency. Inverted1: glass/ITO/PEDOT:PSS/CH3NH3PbI3-xClx/PCBM/TmPyPB/Ag. Inverted2: glass/ITO/Poly-TPD/CH3NH3PbI3/PCBM/LiF/Ag. Normal1: glass/ITO/ALD-TiO2/mp-TiO2/CH3NH3PbI3/Spiro/gold. Normal2: glass/ITO/ SnO2/MgO/CH3NH3PbI3/Spiro/gold. Normal3: glass/FTO/c-TiO2/mp-TiO2/CH3NH3PbI3/Spiro/gold. Flexible1: PET/ITO/SnO2/mp-TiO2/ CH3NH3PbI3/Spiro/gold. Carbon1: glass/FTO/c-TiO2/mp-TiO2/mp-ZrO2/mp-carbon (CH3NH3PbI3). a Extracted from the other reported values. b Linearly scaled from the reported data. c Averaged value.

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on inorganic materials (Si, IIIeV), the literature on this topic is not extensive and further studies are required to better understand the operation and physics of PPV devices under low light environments. In particular, the roles played by the different layers, interfaces, and architectures with a reduced density of photogenerated carriers need further investigation. However, the existing literature can already give some clues for designing efficient PPV cells in low-light environments, which will be discussed in this section. 10.5.4.1 Role of interfacial trap states It is largely accepted that trap states at interfaces and grain boundaries have an important impact on the one sun performance of PPV, as they are known to induce nonradiative recombination [56]. In low-light environments, due to the reduced amount of photogenerated charges, it is believed that these traps may play an even more critical role [17]. To reduce the trap density at the perovskite/ETL interface of iPPV, Chen et al. modified the ETL material and fabrication process [17]. They deposited either C60, C70, or PCBM on top of the annealed perovskite layer. For PCBM, it was deposited by using either a one-step or a two-step deposition process. For the two-step process, a highly diluted solution of PCBM was first spin-coated onto the surface of the perovskite film and annealed at 100 C. Following this, a more concentrated PCBM solution was spincoated. The two-step PCBM ETL resulted in a higher PCE under fluorescent lamp illumination (at 100, 600, and 1000 lux) compared to the other ETLs. Transient photocurrent measurements indicated a more efficient carrier sweep-out and carrier collection for the two-step PCBM, which was explained by better trap filling at the perovskite surface and grain boundaries. Similarly, Li et al. modified the surface of the PCBM ETL by using an ionic liquid of 1-butyl-3-methylimidazolium tetrafluoroborate ([BMIM]BF4) [56]. They showed that the incorporation of the ionic liquid could facilitate electron extraction at this interface and passivate trap states, resulting in Pmax of 98 mW/cm2 under fluorescent lamps of 1000 lux (Fig. 10.8A and B). In another study, Dagar et al. incorporated a thin MgO interfacial layer in between the SnO2 ETL and the perovskite film in a standard device architecture [19]. When measured under indoor white LED illumination, the Pmax was 20.2 mW/cm2 and 41.6 mW/cm2 at 200 and 400 lux, respectively (Fig. 10.8C). It was shown that the thin MgO film can reduce interfacial carrier recombination, which also results in much less hysteresis under low-light illumination and better shelf life stability.

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Figure 10.8 (A) JeV curves of champion perovskite photovoltaic (PPV) (9 mm2) based on PCBM, PCBM/BCP, and PCBM/[BMIM] BF4 ETLs under 1000 lux indoor irradiation. (B) A schematic of the different passivation effects by BCP and [BMIM]BF4, (C) JeV characteristics of PPVs based on SnO2 and SnO2/MgO composite ETLs under 200 and 400 lux white LED irradiation, (DeF) Device properties of PPVs using different architectures and interlayers: (D) Dark JeV characteristics, (E) correlation of the ratios of Pmax at 1000 lux/ Pmax at one sun to series resistance, and (F) VOC measured at different light intensity under AM1.5G illumination and ideality factor calculated. ETL, electron transporting layer; LED, light-emitting diode; PCE, power conversion efficiency. (B) Reproduced from M. Li et al., Interface modification by ionic liquid: a promising candidate for indoor light harvesting and stability improvement of planar perovskite solar cells, Adv. Energy Mater., 8 (2018) 1801509, with permission from John Wiley and Sons) (C) Reproduced from J. Dagar, S. Castro-Hermosa, G. Lucarelli, F. Cacialli, T.M. Brown, Highly efficient perovskite solar cells for light harvesting under indoor illumination via solution processed SnO2/MgO composite electron transport layers, Nanomater. Energy, 49 (December 2017) (2018) 290e299, 2018 with permission from Elsevier) (F) Reproduced from H.K.H. Lee et al., Outstanding indoor performance of perovskite photovoltaic cells e effect of device architectures and interlayers, Solar RRL, (3) (2018) 1800207 with permission from John Wiley and Sons.

10.5.4.2 Importance of cell architecture and interlayers To better understand the roles played by device architectures and interlayers under indoor lighting, we have studied PPV cells under fluorescent lamp illumination with three benchmark device architectures e mPPV, iPPV, and cPPV, as described in Fig. 10.7 [21]. Additionally, alternative HTLs are investigated for the mPPV (using Spiro-OMeTAD or PTAA HTL) and iPPV cells (using PEDOT:PSS or Poly-TPD HTL). These cells show substantial differences in their performance depending on the choice of interlayer, and their performance under one sun is not representative of their

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performance under indoor lighting. The mPPV cell with Spiro-OMeTAD HTL shows a Pmax of >111 mW/cm2 under 1000 lux, while the Pmax of the mPPV cell using a PTAA HTL drops to almost zero under indoor lighting despite these cells have similar performance under one sun. On the other hand, when commonly used PEDOT:PSS is replaced by Poly-TPD as the HTL in the iPPV cell, the Pmax under indoor lighting improves significantly. The performance under low-light intensity correlates well with the leakage current, as shown in Fig. 10.8D. Indeed, if the leakage current level is high, it has a stronger effect on the collection of the photogenerated current under low light conditions than under higher light conditions, as the photogenerated current under low light is smaller. A leakage current density 1 mA/ cm2 is recommended to avoid influencing the performance at low light intensity. Just by changing the HTL, the leakage current can vary by three orders of magnitude, which leads to dramatic difference in PCE under low light even if the PCE under one sun is similar. Brown et al. showed similar results [18]: by using an atomic layer deposited TiO2 ETL (ALD-TiO2) in a standard PPV cell, they obtained much lower dark current than in cells with TiO2 deposited by other coating methods (spray pyrolysis and sol-gel) because of reduced pinhole formation. This resulted in good performance at 200 lux for the ALD-TiO2 cells, whereas the other cells performed poorly. On the contrary to the leakage current, the RS of PPV devices, which plays a major role under one sun AM1.5G illumination, is less critical at low light levels. In Fig. 10.8E, we plot the ratio of the Pmax at 1000 lux to the Pmax at one sun (Pmax,1000lux/Pmax,1sun), along with the RS measured for PPV cells with different architectures and HTLs. A clear correlation between the PCE ratio and the RS can be observed, with increasing values of Pmax,1000lux/Pmax,1sun for higher RS. This demonstrates that a high RS is more critical for the device performance under one sun than under low light levels [14]. In other words, there is a higher tolerance to the RS when the operating current is much lower. This explains why the cPPV cell, which typically shows lower performance under one sun AM1.5G due to a high RS, performs well under low light levels. As described previously, the VOC is a key parameter for PV cells operating under low light intensity. In Fig. 10.8F, the VOC is plotted as a function of light intensity under AM1.5G illumination [21]. By fitting the VOC data to Eq. (10.1), the diode ideality factor n for PPV cells with different architectures and interlayers was calculated. All the cells have diode ideality factor between 1 and 2, except the mPPV-PTAA cell whose high ideality

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factor is affected by the high leakage current in this case. These results show that to achieve a high VOC at low light intensity, the VOC at one sun is more significant than the diode ideality factor, as long as devices possess a reasonable value of n lying between 1 and 2. It is also shown that the VOC could be enhanced by interfacial engineering, e.g., just by replacing the PEDOT:PSS layer with Poly-TPD in the iPPV devices, the VOC was enhanced by 0.120 V at one sun, 0.063 V at 1000 lux, and 0.453 V at 200 lux (all measured in reverse scan).

10.6 Applications 10.6.1 Consumer products Since the introduction of the solar-powered calculator, Teal Photon, in 1978, a large range of other PV products have been developed in the field of product-integrated photovoltaics (PIPV). In a recent study, 90 PVpowered products were analyzed, and it was found that 14% of these products are used indoors and 19% can be used both indoors and outdoors [1]. It is expected that more PV products with low power consumption will emerge in the following years, such as sensors and microenergy harvesting systems for the IoT. Contrarily to outdoor applications, which consist of mainly high-powered PV products, PV products designed for indoor applications should have low power consumptions, ideally less than a few mW [58]. This is because the indoor irradiance, typically in the range 0.01e1 mW/cm2, is significantly lower than outdoors. Most PIPV products (96%) are autonomous, which means that they are not connected to the grid [59]. Among these autonomous products, some of them use an energy storage device (battery or capacitor) which is charged by the PV cell, while others are directly powered by the PV cell without the use of any battery. There are several reasons why indoor light harvesting is combined with or is used as an alternative to batteries. For example, PV cells could overcome the problem of device life span, which is usually limited by the life span of the battery, especially for devices in which it is impossible or difficult to replace batteries [60]. PV cells could also resolve the problem of self-discharge for devices which remains dormant for extended periods of time. Additionally, when a battery is used to store the energy harvested by the PV cell, a stable voltage can be provided to the product [61]. Matching of the battery voltage with the PV cell is usually done by creating a small PV panel with the right number of PV cells connected in series for an appropriate voltage or, for more advanced systems, by using a DC/DC converter.

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In 2007, a polymer PV cell was used to power a lithium-polymer battery with attractive characteristics such as low weight (