Energy-Sustainable Advanced Materials 3030574911, 9783030574918

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Mark Alston Timothy N. Lambert  Editors

Energy-Sustainable Advanced Materials

Energy-Sustainable Advanced Materials

Mark Alston  •  Timothy N. Lambert Editors

Energy-Sustainable Advanced Materials

Editors Mark Alston Faculty of Engineering University of Nottingham University Park, Nottingham, UK

Timothy N. Lambert Department of Photovoltaics and Materials Technology Sandia National Laboratories Albuquerque, NM, USA

“Springer International Publishing AG (outside the USA)”. ISBN 978-3-030-57491-8    ISBN 978-3-030-57492-5 (eBook) https://doi.org/10.1007/978-3-030-57492-5 © National Technology & Engineering Solutions of Sandia, LLC 2021 This work is subject to copyright. All rights are reserved by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors, and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, expressed or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publisher remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Switzerland AG The registered company address is: Gewerbestrasse 11, 6330 Cham, Switzerland

Preface

ESAM presents a collection of chapters that capture the state of the art, and new leading directions, for the synthesis, assembly, and manufacturing of advanced functional materials and systems based on abundant elements, materials, and commodities for application in sustainable/renewable energy generation, capture, storage, and/or modulation. The chapters integrate the detailed fundamental aspects that enable the technical advancements along with an emphasis on the necessity for highly sustainable materials to enable real impact for the benefit of mankind. The book highlights recent developments and the remaining challenges and obstacles that must be overcome. The book describes these principles through several themes: • Evaluating the state of the art in sustainable long-cycle-life batteries. • Investigating electrochemical stability and electrolyte compatibility. • Applying thermal management control to battery design at high operational temperatures. • Modeling analysis of advanced materials for formulated electronic conductivity. • Defining nanostructuring of hybrid materials for enhanced ion and electronic transport. • Fabricating nanotextured structures for enhanced optical absorption of photocarriers, in preparing optimum antireflection layers and device substrates in solar cells. • Development of new thermal energy storage fluids based on interactions between nanomaterials and molten salts. We hope you enjoy reading as much as we have in developing ESAM for your review. Finally, we offer a sincere thanks to our contributors for capturing the innovation of sustainable advanced materials for energy applications and Springer for their assistance in the book publication. Mark and Tim, v

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Preface

This work was supported by the U.S. Department of Energy, Office of Electricity, and the Laboratory Directed Research and Development program at Sandia National Laboratories. Sandia National Laboratories is a multi-program laboratory managed and operated by National Technology and Engineering Solutions of Sandia, LLC., a wholly owned subsidiary of Honeywell International, Inc., for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-NA-0003525. The views expressed herein do not necessarily represent the views of the U.S. Department of Energy or the United States Government. Dr. Imre Gyuk, Energy Storage Program Manager, Office of Electricity, is thanked for his financial support.

Contents

 Aqueous Mn-Zn and Ni-Zn Batteries for Sustainable Energy Storage������    1 Damon E. Turney, Gautam G. Yadav, Joshua W. Gallaway, Snehal Kolhekar, Jinchao Huang, Michael J. D’Ambrose, and Sanjoy Banerjee  Recent Developments of Zinc-Ion Batteries��������������������������������������������������   27 Jaekook Kim, Vinod Mathew, Balaji Sambandam, Muhammad Hilmy Alfaruqi, and Sungjin Kim Molten Sodium Batteries��������������������������������������������������������������������������������   59 Erik D. Spoerke, Martha M. Gross, Stephen J. Percival, and Leo J. Small  Polymer Nanocomposites for Ion Transport ������������������������������������������������   85 Christina A. Bauer  Efficient Light Harvesting in the Nanotextured Thin Film Solar Cells��������������������������������������������������������������������������������������  129 Mohammad Mahdi Tavakoli  Nanomaterials Enhanced Heat Storage in Molten Salts������������������������������  153 Xiaotong Guo, Di Hu, Linpo Yu, Lan Xia, and George Z. Chen Index������������������������������������������������������������������������������������������������������������������  171

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Aqueous Mn-Zn and Ni-Zn Batteries for Sustainable Energy Storage Damon E. Turney, Gautam G. Yadav, Joshua W. Gallaway, Snehal Kolhekar, Jinchao Huang, Michael J. D’Ambrose, and Sanjoy Banerjee

Abstract  Energy storage is a key hurdle for the transition of electrical systems to sustainable solar and wind power. Massive deployment of solar and wind with energy storage is needed for significant reduction of greenhouse gas emissions, requiring manufacturing scales on par with the global automobile industry. Here we review the materials chemistry of rechargeable aqueous Mn-Zn and Ni-Zn because they have outstanding characteristics of sustainability, cost, and safety. Mn cathodes are found to be evolving rapidly under recent research, to potentially offer a breakthrough in cost, and to remain mostly untested at large scale. The limiting chemical processes are (1) the Zn anode cycle life, and (2) the Zn-ion crossover through the separator, for both of which recent research is presented. Keywords  Batteries · Energy storage · Electrochemistry · Alkaline battery · Zinc · Manganese · MnO2 · Aqueous battery · Grid storage · Sustainable energy · Smart grid · Green chemistry · Clean technology · Clean energy · Renewable energy

1  Background Battery systems for energy storage are a key component to enable rollout of solar and wind power, allowing reduction of greenhouse gas emissions. They also mitigate toxic air emissions such as mercury from coal power or NOx, smog, and particulates from gas-fired peaker plants or internal-combustion vehicles. Sustainability of the battery systems themselves is an emerging research field motivated by rapidly growing markets for large-scale energy storage. Until approximately 2015, the global battery market predominantly served applications below 100 Wh and produced an order of magnitude of ~50 GWh (~1 billion kg) of product per year [1, 2]. At this small scale, the material supply chains did not cause major environmental concern other than recapturing toxic metals from end-of-life nickel-cadmium D. E. Turney (*) · G. G. Yadav · J. W. Gallaway · S. Kolhekar · J. Huang M. J. D’Ambrose · S. Banerjee Energy Institute, Department of Chemical Engineering, City University of New York, New York, NY, USA e-mail: [email protected]; [email protected] © National Technology & Engineering Solutions of Sandia, LLC 2021 M. Alston, T. N. Lambert (eds.), Energy-Sustainable Advanced Materials, https://doi.org/10.1007/978-3-030-57492-5_1

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(Ni-­Cd) and lead-acid batteries. These recycling programs are very successful in most developed nations, but unfortunately are not in many developing countries [3, 4]. Going forward from 2015, the emergence of large-scale applications of 100 kWh to 100 MWh for electric vehicles (EVs) and grid-scale energy storage created predictions that the global market for batteries will exceed 2000 GWh (~40 billion kg) per year by 2030 [5, 6]. This enormous increase in production creates an urgent need to address issues of sustainability. For example, the eventual maximum global market size for batteries will likely top 500+ GWh per annum [7], (Jaffe et al. 2016), with current technology would require ~100,000+ metric tons of lithium per year, an amount three times the current lithium global production rate [8], and which would create overall supply chain material flows on par with the global automobile industry. Therefore, the mining, processing, and recycling of rare elements for future battery production becomes an important issue. Government regulation, life cycle assessments, and recycling strategies should be developed for handling sustainability issues such as greenhouse gas emissions, landscape and mining impact, toxics release, and fire/industrial safety [9]. Battery system cost is analogous to sustainability when full life cycle costs are considered. As of 2019, the cost of greenhouse gas emissions and avoidances are not factored into battery prices, nor are the recycling or end-of-life costs, but, as a rough approximation, battery system cost is a convenient indicator of the resources needed for production, and it is certainly the most important metric for short-term market sustainability. Estimates of the battery system cost needed to achieve profit in common electrical grid applications are given by Eyer and Corey [7], showing costs must be well below $100 per kWh for most industrial or utility applications, and as low as $20 per kWh for transmission congestion relief. The technology development curve for lithium ion is now widely expected to level off near $300 per kWh by 2020 and perhaps to $200 per kWh by 2030 [10, 11], meaning that lithium ion will not be financially positioned to serve most grid applications. Figure 1 shows cost estimates of aqueous Mn-Zn and Ni-Zn [12–15], could reach below $50 per kWh suggesting they are perhaps the most sustainable battery chemistry, at present date of 2019. Mn-Zn batteries also show superior fire safety and toxicity risks, which will be discussed later. Ni-Zn is of interest as a proxy study for Mn-Zn, so will be discussed in this chapter. Zinc-air and zinc-bromine batteries are not considered here due to their lack of demonstrated cycle life at industrial scale or in industrial settings. Aqueous silver oxide (Ag-Zn) and nickel metal hydride (Ni-MH) batteries with very high cycle life and energy density (e.g., [16]) will not be considered here due to the high cost of silver and nickel. An important point made throughout this chapter is the importance of academic researchers to be practical and transparent, to disclose mAh/mL cycled and mAh-lifetime/mL (or per gram) with the full materials used, not a subset. Cost models to accurately show $/kWh and $/ kWh-lifetime are also important. To highlight the great interest in these chemistries in recent years, Fig. 2 shows the intensity of research literature on Mn, Zn, and Ni battery systems.

Aqueous Mn-Zn and Ni-Zn Batteries for Sustainable Energy Storage

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Fig. 1  A market application map of grid energy storage, showing market positions of leading battery technologies. The colored numbers inside each color shaded area denote achievable cycle life at that specific cost per kWh

Fig. 2  History of published research on Zn, Mn, or Ni electrodes, showing spikes in activity due to (a) industrial rollout of modern alkaline Mn-Zn cells, (b) alkaline Mn-Zn overcoming Le Clanche cells in the portable electronics market, (c) Ni-Cd and Ni-MH enter the industrial markets, (d) proliferation of mobile telephones, and (e) growing awareness that energy storage is critical for wind and solar power grid penetration

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2  Present State of Mn-Zn and Ni-Zn Commercial Technology 2.1  Historical Rechargeable Manganese-Zinc Batteries Single-discharge Mn-Zn cells are one of the most commercially successful battery technologies in history, as they comprise the popular products from Energizer, Duracell, Eveready, and other battery market competitors. They were first created by Karl Kordesch, Lewis Urry and coworkers as reported, for example, in US Patent 2,960,558 in 1960. The Mn-Zn chemistry holds ~200 Wh per kg in an off-the-shelf single-use Mn-Zn cell, at a retail cost of $40 per kWh. This outstanding performance motivated attempts to develop rechargeable Mn-Zn as early as the 1970s. Briefly, the chemistry is believed to operate by a reaction pathway similar to

MnO2  H 2 O  e   MnOOH  OH  at ~ 1.1Vvs.Zn 2 



2 MnOOH  H 2 O  2OH   Mn  OH 4

 MnO2  H 2 O voltage independent

MnOOH  H 2 O  e   Mn  OH 2  OH  at ~ 0.9V vs. Zn 2 

2 MnOOH  Zn  OH 4

2 

2 MnOOH  Mn  OH 4

(1)

(2) (3)



 ZnMn 2 O 4  2OH  2H 2 O voltage independent (4)  Mn 3 O 4  2OH  2H 2 O voltage independent



(5)

where (1) and (3) are the energy storage reactions. Only specific crystal isomorphs will yield these reactions, with reaction (1) being best done with “electrolytic MnO2” (EMD) and reactions (4) and (5) creating inactive byproducts hausmannite (Mn3O4) or haeterolite (ZnMn2O4) which are irreversible. Further details are available in references [17–19]. More successful cycle life historically relied on additives such as titanium, bismuth, lead, or barium to reduce the dissolution reaction of 2 MnOOH to form Mn  OH 4 thus preventing the formation of inactive hausmannite [20–22]. Such chemistry formed the commercialization attempts of Rayovac RENEWAL cells in the early 1990s, but unfortunately the capacity of this and similar additives did not typically survive over ~20 cycles, and customers did not accept them. Recent work has continued with similar additives as discussed in Sect. 3.1, e.g., see Table 1. Related developments at Ford Motor Co. in the late 1980s discovered bismuth additives to enable deep-­cycling (close to 2 e− per MnO2, 600 mA g−1) and long cycle life of MnO2 under linear-voltage-sweep cycling conditions [26], but unfortunately not during galvanostatic or constant power cycling, thus preventing commercialization. A later company (RBC) continued research on this technology but has not overcome this problem.

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Aqueous Mn-Zn and Ni-Zn Batteries for Sustainable Energy Storage

Table 1  Performance comparison of the manganese electrodes from recent publications, where n.r. denotes “non-reported” measurements in a publication capacity per capacity per capacity per solidsa volumea capacity per area mass (mAh g−1) (mAh g−1) (mAh cm−3) Reference (mAh cm−2) Cycle life Wroblowa and 17 521 n.r. n.r. 20 Gupta [106] Kannan et al. n.r. ~375 ~187 n.r. 600 [107] 210 147 n.r. 30 Raghuveer and n.r. Manthiram [108] Ingale et al. [27] 5 31 15 50 4000 Hertzberg et al. n.r. ~250 ~125 ~100 50 [109] Hertzberg et al. 26 ~425 ~255 ~200 60 [110] Rus et al. [111] n.r. 225 147 ~110 30 Pan et al. [23] 1.4 280 196 ~140 5000 Yadav et al. [14] 29 617 370 315 3500 Zhang et al. [24] 2.5 225 ~200 ~110 2000 “Total solids” and “total volume” are for the manganese electrode and pore-electrolyte only, and do not include the mass or volume of the Zn anode electrode b Also see Kordesch et al. [25] a

On the zinc side, the end products of a still-unknown reaction pathway require the overall stoichiometry to follow 2 



Zn  4OH   Zn  OH 4 2 



Zn  OH 4

 2e  at 0.0 Vvs.Zn

(6)



 ZnO  H 2 O  2OH  voltage independent



(7)

where in the failure mechanism is usually reported as “shape change” in which the zinc balled up in isolated locations and refused to cycle. The stringent economics of grid-scale energy storage became important during the US Department of Energy’s ARPA-E GRIDS program of 2009–2013. In that program a pathway was discovered via Mn-Zn technology for battery system cost to be reduced to nearly $100 per kWh and maintain long cycle life. This technology employed a Mn cathode in which only 0.1–0.3 electrons cycle per Mn atom (EMD) and only 0.1 electrons cycled per Zn atoms. Cycle life was 1500–4000 and costs were below ~$200 per kWh energy storage [27, 28]. This Mn-Zn technology could replace Pb-acid batteries and eliminate significant environmental pollution from lead. Also, life cycle assessment finds the greenhouse gas emissions from Mn-Zn technology to be better than Ni-Zn or lead-acid batteries if the Mn is cycled at more than 20% of the first electron of MnO2 [29]. Commercialization of such technology is under develop-

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ment at Urban Electric Power corporation as of 2020. EMD can theoretically cycle up to 0.7 electrons per Mn [19] if the MnOOH material can reversibly intercalate H+ into EMD’s mixture of pyrolusite and ramsdellite crystals, as shown in Fig. 3. The effect of depth of zinc discharge in MnO2-zinc batteries on the overall cycle life is shown in Fig. 4. A separate study by [28] cycled EMD Mn-Zn cells at 0.1 electron per Mn and

Fig. 3  A conceptual picture of the Mn rechargeable battery system, which is stable cycling 0.0 to 0.79 electrons per MnO2, or can achieves 2.0 electrons per Mn when the Cu-Bi-birnessite system of Yadav et al. [14, 15] is employed

Fig. 4  Cycle life as a function of zinc depth-of-discharge for shallow-cycled manganese-zinc batteries. A “S1” cell is a ~ 1 L Mn-Zn cell from Urban Electric Power. The red and gray asterisks are outlier experiments that failed due to anomalous events

Aqueous Mn-Zn and Ni-Zn Batteries for Sustainable Energy Storage

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found 30% of the cells failed due to zinc short-circuits, 56% failed due to the Zn electrode passivating shape change, 14% were purposefully destroyed for dissection, and 0% failed due to problems with the Mn cathode. The bismuth-doped rechargeable Mn cathode pioneered in the 1980s by Wroblowa et al. created hope for a fully rechargeable Mn cathode in alkaline electrolyte. Wroblowa et al. used Bi2O3 additive to enable Mn cycling at 2 electrons per Mn without formation of the inactive phases such as Mn3O4, but they were only able to do so with linear-voltage-sweep cycling, and could not achieve the same success with galvanostatic cycling for unknown reasons. Full rechargeability of Mn at two electrons per Mn atom has therefore been of renewed research interest and will be discussed in Sect. 3.1.

2.2  Historical Nickel-Zinc Batteries Significant markets exists for large-format high-power (>50 W L−1) batteries that can deliver energy density (>100 Wh L−1), inherent fire safety, low maintenance, and preferably a 5+ year lifespan and cycle life over 300 depending on the application. This ratio of power capacity to energy capacity is more than most lead-acid batteries can provide [16]. Although lithium-ion technology dominates this market niche for small scale applications, its many thermal runaway accidents have significantly delayed its penetration into fire-sensitive environments, e.g., aircraft electrical systems [30, 31]. In such situations, Ni-MH and Ni-Cd today are preferred over lithium ion. Industrial uses of Ni cathodes are highly successful. As one example, Ni-MH batteries in hybrid-vehicles such as the Prius now have warranties over 150,000 vehicle miles, and NYC taxi vehicles with Ni-MH regularly achieve 300,000 miles without battery replacement [32]. Given the success of Ni-MH and Ni-Cd, there has long been motivation to develop Ni-Zn batteries due to their 1.6 V output as compared to 1.2 V for Ni-MH, and also due to the lower cost of Zn compared to metal hydride, and non-toxicity compared to cadmium. Many corporations have attempted commercialization of Ni-Zn, beginning with Edison’s work in 1901, and with recent attempts by Evercel, PowerGenix, Evionyx, ZincFive, ZAF, Enersys, and BASF. The electrochemistry of Ni-Zn is summarized by

NiOOH  H 2 O  e   Ni  OH 2  OH 



Zn  4OH   Zn  OH 4  2e  2





(8) (9)

where fuller details exist in other literature [12, 16, 33, 34] and for the Zn electrode [28, 35]. Briefly, the Ni material cycles between β-Ni(OH)2 and β-NiOOH, both of which are layered nickel hydroxide sheets with protons filling the interlayer ­galleries, as shown in Fig. 5. The β-Ni(OH)2 is the reduced form, and is a good ionic

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Fig. 5  Crystal structures are shown for Ni(OH)2 on the left and NiOOH on the right. Red atoms are oxygen, gray are nickel, and small spheres are hydrogen

conductor but poor electrical conductor. The β-NiOOH material is oxidized sufficiently to react with H2O, thus the β-Ni(OH)2 is thought to coat the β-NiOOH and prevent it from self-discharging by contact with H2O. A few percent Co(OH)2 and often Cd(OH)2 is usually mixed with the NiOOH to improve cycle life and reduce oxygen evolution reaction [36]. The zinc electrode in a Ni-Zn is usually the first to fail, and a very long list of additives and methods are researched to solve the problem [35], with the most successful being addition of Ca(OH)2 [28, 37] and perhaps with developments from PowerGenix corporation [38]. One recent advance is the development of “flow-­ assist” Ni-Zn wherein the zinc side is flowing electrolyte, which will be discussed in Sect. 3.2. A demonstration project funded by the US National Energy Technology Laboratory and ConEdison of New York showed a 30 kWh “flow-assist” Ni-Zn battery to cycle at 1100 healthy cycles at 2 h rates and ~ 75% depth-of-discharge of the available zinc [13], which demonstrated a cost estimate for manufacturing this technology (including US labor) of $400 per kWh. An example of a commercial Ni-Zn battery recently produced by Evercel is shown in Fig. 6. Energy capacity of this Evercel Ni-Zn technology circa 2001 was 110  Wh/L, 65  Wh/kg, with cycle life near 500 when one-hour charge–discharge times are used and 20% of theoretical capacity is cycled. Sintered nickel plaque provides the best cycle life, but due to cost concerns the Ni side is often a carbon-­ paste PTFE-bound NiOOH electrode (US Patent 4,546,058).

2.3  Separators in Use for Alkaline Batteries Separators command a large fraction of total battery cost, up to 25% in some cases [39], and play a key role in performance. Extensive work has been performed on separators for lithium-ion batteries for preventing catastrophic short-circuits and electrolyte combustion [40]. Aqueous batteries usually have no risk of thermal run-

Aqueous Mn-Zn and Ni-Zn Batteries for Sustainable Energy Storage

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Fig. 6  A commercial Ni-Zn battery product from Evercel corporation, circa 2001. The pasteNiOOH (left) and paste-ZnO (right) electrodes used in this technology are shown to the right side of the battery

away, thus less research attention has accumulated on aqueous separators. Manufacturer’s cost appetite for separators in aqueous systems is less than in other battery categories due to overall system cost being lower, but the choice of separator has profound effect on cycle life of Zn battery systems due to protection from zinc dendrites that may short-circuit, and also due to protection from zinc electrode shape change. To prevent short-circuits a nano-pore cellophane membrane is found in almost all zinc-alkaline battery systems, including some single-use cells. Celgard or similar polymer layers with sub-micron pore size are sometimes substituted for the cellophane but cost often limits this substitution. In addition to the cellophane layer, a water-absorbent fleece layer (a.k.a. non-­ woven felt) is also commonly placed next to the zinc electrode to prevent drying out. Pellon is a proprietary version of this water-absorbent layer and is used when chemical stability of the separator is an issue. Several variants of Pellon are offered commercially. An approximate listing of the various wicking layers for alkaline systems using a zinc anode is given in Arora and Zhang [40].

3  Next-Generation Materials for Mn-Zn and Ni-Zn Systems 3.1  Advanced Rechargeable Manganese Cathodes MnO2 chemistries that cycle ~2 electrons per Mn were recently published for both alkaline and acidic electrolytes [14, 15, 23, 41]. Pan et al. [23] reported long cycle life of ~285  mAh/g from a cathode of α-MnO2 pressed with carbon powder and

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infiltrated with aqueous 2 M ZnSO4 electrolyte. The success of Pan et al. was due to additive manganese sulfate (0.1 M MnSO4) in the electrolyte, which prevented loss of Mn from the electrode during the energy storage reactions. This same chemical method and experimental cyclability was reported by a group of Korean researchers in 2001 (US Patent US6187475 B1). Unfortunately, this acidic-MnO2 cathode worked only with a small loading of active material (1.4 mAh/cm2), which will keep overall battery costs high. It is promising that Pan et al. pair the α-MnO2 cathode with a Zn foil to claim a new battery with ~170 Wh kg−1 (where the mass in their calculation ignores the non-cycling Zn foil). Another recent breakthrough succeeded to cycle ~2 electrons per Mn (617 mAh/g) for 6000 full cycles with practical loading of active materials (30 mAh/cm2) by adding bismuth and copper to EMD or birnessite δ-MnO2 ([14, 15], 2018). Only previous literature from Wroblowa’s group [42, 43] cycled the complete 617  mAh/g Mn capacity reversibly, but they were unable to cycle even half that capacity galvanostatically. Yadav et al. [15, 44, 45] showed full-cell cycling of a Zn-Mn battery at 2-electrons per Mn and high Zn utilization, with chemistry as shown in Fig.  7, reaching several 100 cycles and energy density above 100 Wh/L based transparently upon full cell specifications. This new Mn-Zn technology is appropriate for grid-scale energy storage as shown in Fig. 1, and perhaps even for transportation uses (Fig. 7). Table 1 shows a survey of performance for recent work on manganese electrodes. The key metrics for comparing the practicality of the Mn cathodes are cycled capacity per area (mAh  cm−2), cycled capacity per total cathode volume (mAh  cm−3), capacity per total solids mass (mAh  g−1), and cycle life. Combining the Cu-Bi-­ birnessite cathodes with a zinc anode creates an impressive cell with the best performance shown in Fig. 1. Unfortunately the zinc ions have a poisoning effect on the birnessite-MnO2 [14, 15, 43] by causing the formation of haeterolite ZnMn2O4 [15]. Most common separators allow the zinc to cross from the anode to the cathode material, thus zinc-blocking separators are the focus of Sect. 3.4. Further, the zinc electrode suffers from poor cycle life when their percent utilization is pushed above ~5%, thus zinc cyclability is the focus of Sect. 3.3. Nevertheless, [15] showed separators of Ca(OH)2 could temporarily block zinc crossover and allow the Mn-Zn cell to cycle for 1000 cycles with an energy density near 160 Wh L−1 based transparently upon all materials used. The relationship between cell energy density, percent loading of Mn active material (for a 0.5 mm cathode with ~0.1 g cm−2 of total mass), and zinc utilization is shown in Fig.  8. Cost of energy storage ($/kWh) and ($/kWh-­ throughput) is very sensitive to the metrics in Table 1. Technologies from [14, 15] and [14, 15] achieve costs under $150/kWh and $50/kWh.

3.2  Current Research on Rechargeable Zinc Anodes As mentioned in Sect. 2.1, 2.2, and 3.2, the zinc electrode is the least reliable component of rechargeable Mn-Zn or Ni-Zn batteries, and usually leaves ~80% of its theoretical capacity un-utilized. This produces great opportunity for improvement, and is perhaps why research on zinc electrodes is still very active. A literature survey

Fig. 7  Conceptual electrochemical understanding of the Cu-Bi-birnessite chemistry when paired with a zinc anode is shown

Fig. 8  A map of total cell energy as a function of DOD of zinc’s theoretical capacity and the loading of MnO2 into the cathode mix (assuming a 0.5 mm cathode with ~0.1 g cm−2 of total mass)

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Table 2  A survey of rechargeable Zn anode performance since the 1980s

Materials comprising Ref. the anode Biegler et al. ZnO, acetylene black, [46] 2% HgO, PTFE, 1% CMC, 6.9 M KOH Sato et al. [47] ZnO, Zn, Bi2O3, Ca(OH)2, Teflon binder Nichols et al. ZnO, PbO, PTFE, 15% [48] KOH w/ KF or K3BO3 w/ 1% Li2BO3 Gagnon [49] ZnO, Ca(OH)2, 3.4%Pb3O4, 1.4% newsprint, 20% KOH Jones [50] ZnO, Ca(OH)2, 3.9%Pb3O4, 1.1% newsprint, 30% KOH Gagnon and ZnO; Ca(OH)2; 3.3% Wang [51] Pb3O4; 1.1% superwettable PP, 20% KOH Charkey [52]) ZnO; Ca(OH)2, PbO, 2.5% PTFE binder Müller et al. ZnO, PTFE, cellulose [53] fibers Adler et al. [54] ZnO, PbO, PTFE, newsprint, 4 M KOH, 2 M KF, 2 M K2CO3, LiOH Charkey [55] Calcium zincate, 50% ZnO, 10% Ca(OH)2 Yu et al. [37] Calcium zincate, 8% PTFE; 2% PbO; 4 M KOH sat’d w/ ZnO Zhang et al. ZnO, 27% Ca(OH)2, [56] 10% Bi2O3, ~1% PVA binder Wang et al. [57] Ca-Zn; Zn powder; acetylene black, PTFE, 6 M KOH sat’d ZnO KOH, 0.75 g/L zincate, Ito et al. [58] electrodeposition from and Turney electrolyte et al. [13] Gan et al. [59] ZnO coated in polypyrrole, 10% acetylene black, PVA, KOH

cycle lifea ~50

DOD: per total anode massb(mAh g−1) 112

DOD: per total anode volumec(mAh mL−1) 338

lifetimec throughput (Ah mL−1) 17

300

128

330

99

130

142

242

31

200

123

291

58

300

125

364

110

500

91

225

112

450

218

396

178

175

n.d.

246

43

380

170

340

131

170

164

186

32

400

174

428

160

100

127d

498d

50d

230

240

390

88

3500

125

29

103

70

200

588

41

(continued)

Aqueous Mn-Zn and Ni-Zn Batteries for Sustainable Energy Storage

13

Table 2 (continued)

Ref. Chamoun [60] and Davies et al. [61] Ingale et al. [27] Wang et al. [57] and Zhang et al. [62] Parker et al. [63] Higashi [45] Turney et al.[28]

Materials comprising the anode Hyper-dendritic zinc, ~6 M KOH

cycle lifea ~100

DOD: per total anode massb(mAh g−1) n.r.

Zn, ZnO, PTFE

2000

11

38

72

ZnAl-X-LDH, acetylene black, PTFE, 6 M KOH sat’d ZnO, Zinc sponge created by electrodeposition “Backside” cycling ZnO, Ca(OH)2, Bi2O3, Teflon binder, KOH

800

200

242

193

>40

~100e

~600

>24e

800 990

~50f 58

31f 192

25f 192

DOD: per total anode volumec(mAh mL−1) n.r.

lifetimec throughput (Ah mL−1) n.r.

Cycle life was defined as the number of cycles until capacity drops below 75% of the average capacity, or until coulombic efficiency drops below 75%. Many publications that reported low capacity (1000  cycles) in 1  M ZnSO4 at very high current rates of 15C (1C = 300 mAh g−1) was demonstrated [62]. In situ XRD analysis predicts the role of water molecules, which is crucial for reversibly expanding and contracting the layered galleries to allow Zn2+ ingress/egress, leading to good kinetics and high rate performance. It is evident that Zn2+/H2O pillared layered V2O5 and the effect of water molecules (both lattice and/or water molecules from the electrolyte) expand the layered galleries and buffer the high charge density of the intercalating ions; thereby electrochemical reactions are greatly enhanced. A schematic representation of Zn-ions intercalation/de-intercalation is represented in Zn0.25V2O5·nH2O gallery, as shown in Fig. 3a. The cycling results, in Fig. 3b, indicate that the corresponding cell can be cycled for over 200 cycles at 1200 mA g−1 current density. An ease of overall battery assembly/fabrication with an aqueous electrolyte (1  M ZnSO4), a metallic Zn negative electrode and scalable processing of the Zn2+ host material, help in meeting the requirements essential for large-scale storage applications. Hence, many layered and tunnel-type vanadium oxides including LiV3O8 [63], VO1.52(OH)0.77 [64], Zn3V2O7(OH)2·2H2O [65], H2V3O8 [66], V2O5·nH2O [67], Na0.33V2O5 [68], Na1.1V3O7.9 [69], and Zn2V2O7 [70] have been investigated recently as potential cathode materials for ZIBs. Similar to the case of manganese-based electrodes, the effect of different metal ions in the interlayers, which act as pillars to not only holding the backbone and increasing the stability of the vanadium oxides but also intensively improve the ion diffusion rate, as well as the electronic conductivity, have been studied for ZIBs. Recently, Na0.33V2O5 nanowire cathode delivered a very high capacity of 367 mAh g−1

38

c

Charge cycling

5

10

15

19

25

30

35

1

2

3

4

5

4

13.5

d

15.0 2θ /°

(100)

1

5 1

5

10

10

19

20

25

30

35

38

15

21 24 27 2θ /°, λ = 1.5414 Å

e

0

50

100

150

200

250

15

19

20

25

30

35

38

Layered Zn0.25+x V2O5·zH2O

Intensity / a.u.

H2O

Capacity (mAhg-1) 30

0

f )

32 34 36 38 2θ /°, λ = 1.5414 Å

1

5

10

15

19

20

25

30

35

38

40

100 150 Cycle number

Zn-inserted LVO phase

50

42 44 2θ /°

ss

92

1

5

10

15

19

20

25

30

35

38

90 200

Discharge Charge

94

96

98

100 Coulombic efficiency (%)

Fig. 3 (a) Schematic representation of electrochemical intercalation/de-intercalation of Zn-ions into the vanadium oxide bronze, Zn0.25V2O5·nH2O gallery and (b) corresponding cyclability profile over 200 cycles under 1200 mA g−1 (4C). (reprinted with permission from ref. [62]) . (c) Electrochemical discharge/charge profile for LiV3O8 cycled within 1.2–0.6 V. Corresponding in situ synchrotron XRD scans with selected regions (d) 13.2–18°, (e) 20–29.5°, (f) 30–9.5°, and (g) 39.5–45° during the electrochemical reaction. (reprinted with permission from ref. [63])

1 0 1.2 0.9 0.6 Potential/V

Discharge cycling

Zn2+

(003)

Zn (s)

Time/x 10 sec Intensity / a.u.

−H2O

(011)

300

Intensity / a.u.

2+ +Zn /2e-

(.111)

350

03

(1

400

Intensity / a.u.

b

(.301)

Layered Zn0.25V2O5·yH2O

(203)

a

Recent Developments of Zinc-Ion Batteries 39

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J. Kim et al.

at 0.1 A g g−1 using a 3 M Zn(CF3SO3)2 electrolyte and metallic Zn as an anode within a wide potential range of 0.2–1.6 V [68]. However, the electrolyte is seemingly very expensive (roughly 15 times costlier than ZnSO4). These nanowires display a long term of cycle stability with a capacity retention of nearly 93% for 1000  cycles at 1  A g−1 current density. The structural evolutions during electrochemical reaction were studied using XRD and Raman and X-ray photoelectron spectroscopies. It is found that the indigenous ions can act as pillars to stabilize the layered structure, thereby ensuring an enhanced cycling stability. Surprisingly, based on Zn0.25V2O5·nH2O [62], yet another pillar-stabilized cathode, Ca0.25V2O5·nH2O, a double layered calcium vanadium oxide bronze, has been reported for ZIBs [71]. The V2O5 layers stack along the c-axis and the intercalated metal ions (Zn2+ or Ca2+) as well as water molecules reside in the interlayer space and strongly expand the interlayers. The nanobelts cathode delivered a high capacity of 340 mAh g−1 at 0.2C, good rate capability and prolonged cycling lifespan of 3000 cycles, holding 96% capacity retention at a very high applied current rate of 80C in the potential range of 0.2–1.6 V using Zn(CF3SO3)2 electrolyte. Further investigations on the crucial role of water molecules facilitating structural stability and superior electrochemical performance in layered-type V2O5·nH2O cathodes of ZIBs was also performed [67]. The H2O-solvated Zn2+ possesses largely reduced (screening) effective charge and thus greatly suppresses the electrostatic interactions with the V2O5 framework and thereby effectively promotes its diffusion. The influence of the “lubricating” effect of water molecules thus contribute to the aqueous Zn battery showing an energy density of 90 Wh kg−1 at a high-power density of 6.4 kW kg−1 (based on the active mass of cathode and anode materials). The resulting performance thus make such “structural water-contained” layered oxides a promising candidate for high-performance, safe, and environment-friendly energy storage devices. Layered LiV3O8 (LVO), a well-known candidate for LIBs, follows the monoclinic system which comprises of the corner sharing of two edge shared octahedral (VO6) and trigonal bipyramid (VO5) units to form (V3O8)− layers along the (100) plane, the layers being linked by Li+ ions in the interstitial octahedral and tetrahedral sites. Hence, it is highly feasible that the Zn-ions can well intercalate into the host cathode since both Li+ and Zn2+ have almost similar ionic radii. Alfaruqi et al. demonstrated that the flake-type LVO cathode delivers an average discharge capacity of 172 mAh g−1 at 133 mA g−1 current density after 65 cycles within the potential range of 1.2–0.6 V [63]. Apart from that, the electrochemical regulation through the operando in situ XRD study is a quite interesting and complex phenomenon. The initial stage of Zn-ions intercalation (during discharge) is inferred to the single-­ phase reaction of LiV3O8 where the Zn-ions begin to occupy vacant lithium sites with different energies to form a ZnLiV3O8 phase. In the intermediate stage, the peak splitting feature of the (100) plane appears to suggest a two-phase reaction of LVO and ZnLiV3O8. A complete single-phase reaction to form ZnyLiV3O8, y > 1 at the end of the discharge is noted at the final stage, as shown in Figs. 3c–g. On the other hand, a complete one-step early stage of the electrochemical charging/Zn de-intercalation is represented by the single-phase reaction of ZnyLiV3O8

Recent Developments of Zinc-Ion Batteries

41

that gradually proceeds towards the recovery of LiV3O8. The complete electrochemical reaction through intercalation/de-intercalation is well supported by XRD simulation. Thus, although the layered-type structure of LVO undergoes slight modification during the electrochemical reaction with Zn2+ or Li+ ions, the phase behavior is different from that observed during lithiation. Therefore, the present study encourages the utilization of such an approach to understand the phase evolution in layered-type systems but also demonstrates vanadium-based intercalation hosts as a promising cathode for ZIBs. A similar anionic [V3O8]− layered network containing H2V3O8 nanowire cathode, in which hydrogen atoms are linked between the layers, and each layer consists of VO6 octahedra and VO5 trigonal bipyramids, is utilized for ZIBs. This cathode nanowire of Zn-H2V3O8 battery delivers a high capacity of ~424  mAh  g−1 at 0.1  mA  g−1 within the potential window of 0.2–1.6  V using a 3  M Zn(CF3SO3)2 electrolyte [66]. Interestingly, the nanowire cathode retains nearly 94% of the capacity after 1000 cycles at high current rate of 5 A g−1. In general, the metal ions in [V3O8] layers containing MxVnOm, (where M = metal ion) network can stabilize the structural integrity. A highly durable Na2V6O16·1.63H2O nanowire cathode and an aqueous Zn(CF3SO3)2 electrolyte have been developed [72]. This aqueous Zn-ion battery provides a high capacity of 352 mAh g−1 at 50 mA g−1 and exhibits a capacity retention of 90% over 6000 cycles at 5 A g−1 current density within the potential range of 0.2–1.6  V.  A similar [V3O8]− layered background containing pilotaxitic Na1.1V3O7.9 nanoribbons/graphene cathode shows a reversible capacity of 171 mA h g−1 after 100 cycles at 300 mA g−1 current drain [69]. The results indicate that the micro-structure is stable after long-term cycling for ZIBs. Recently Kundu et  al. reported V3O7•H2O cathode, a similar two-dimensional [V3O8] layers show very high capacity and power (375 mAh g−1 at 1C rate and 275 mAh g−1 at 8C rate) in an aqueous electrolyte of 1 M ZnSO4 [73]. The structural evolution during electrochemical cycling was monitored through operando in situ analysis. Interestingly, a formation/dissolution of an additional new phase, Zn4(OH)6SO4•5H2O evolved during discharge/charge processes, respectively. Though the observation is not quite new, its role on electrochemical reaction is not well defined. According to them, this observation is a side reaction in which dissolved oxygen from the electrolyte may interact with available species (Zn2+, SO42− and H2O) to precipitate on the surface of the electrode. However, further systematic studies are necessary to identify the real contribution in conjunction with redox electrochemical kinetics. Besides the (V3O8)− layered background, another framework of pyrovanadate, (V2O7)4− can stabilize a larger structure as Zn-pyrovanadate, α-Zn2V2O7. This structure consists of a layered structure of tetrahedrally coordinated VO4 and distorted trigonal bipyramid ZnO5 polyhedra. Sambandam et al. reported on an α-Zn2V2O7 nanowire cathode that delivered high specific reversible specific capacities of 197 and 138  mAh  g−1 at two different current densities of 300 and 4000  mA  g−1, respectively, for 200 and 1000  cycles through an insertion mechanism [70]. ­ Significant specific energies of 166.3 and 114.3 Wh kg−1 at specific powers of 36 and 3168 W kg−1 (based on cathode material mass), respectively, were delivered by this cathode using 1  M ZnSO4 aqueous electrolyte within the potential range of

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0.4–1.4 V. In specific, the electrode displays the energy density of 87.3 Wh kg−1 for a given power density of 18.9 W kg−1 at 50 mA g−1 current density (based on whole cell). The lowest energy density of 60 Wh kg−1 is higher than those of commercial Pb-acid (30  Wh  kg−1) and Ni–Cd (50  Wh  kg−1) batteries. Another pyrovanadate, Zn3V2O7 (OH)2·2H2O undergoes structural evolution during electrochemical reaction through intercalation/de-intercalation mechanism [65]. The cathode delivers a long cyclability with reversible capacity of 101 mAh g−1 after 300 cycles with 68% retention using 1 M ZnSO4 electrolyte cycled between 0.2 and1.8 V. VS2 is one among the member of transition-metal dichalcogenides (TMDs), with hexagonal system, which shows similar crystal structure to that of graphite lamellar with an interlayer spacing of 5.76 Å. The aqueous Zn-VS2-nanosheets battery, the first chalcogenide based first report for ZIB, delivers a reversible capacity of 110.9 mAh g−1 at 0.5 A g−1 after 200 cycles with nearly 98% capacity retention in a 1 M ZnSO4 electrolyte solution [74]. The overall capacity in this study is governed by the combination of diffusive-controlled reaction and surface controlled capacitive reaction through cyclic voltammetry study by applying different scan rates. Nearly 62% of the total capacity contribution was confirmed to be from surface capacitive capacity at 0.2  mV  S−1 scan rate. The NASICON-type phosphate Na3V2(PO4)3 (NVP) is a promising cathode and the Zn//0.5 M Zn(CH3COOH)2 // Na3V2(PO4)3 battery configuration showed a reversible capacity of 97 mAh g−1 with 74% capacity retention after 100 cycles at 0.5 °C [75]. Before Zn insertion, initially 2 moles of Na were electrochemically de-intercalated from the NVP structure and reversible Zn-intercalation occurred from the subsequent cycles and the formation of a new inserted phase, ZnxNaV2(PO4)3 was observed. These interesting studies clearly open the door towards the development of safe, eco-friendly, and economically viable electrodes for ZIB applications. The number of published reports for vanadium cathodes based for aqueous ZIBs are quite high, in the last couple of years, thereby making it tedious to report all these results. Table 2 presents a few of the selected works with detailed electrochemical properties including rate performance, concentration of the electrolytes, and additives, if any. Open-Framework Structures There are some special categories of open-framework structures like Prussian blue analogues (PBAs) that have garnered significant attention as hosts for multivalent cations. Zhang et al. reported on an open 3-D framework structure of Zn3[Fe(CN)6]2 (ZnHCF) that was cycled between 0.8 and 2  V and this cathode exhibited good stability with a capacity retention of 76% after 100 cycles when being fully charged/ discharged in 2 h (1C, where 1C = 60 mA g−1) and 81% capacity retention after 100 cycles at a high current drain of 24 min (5C) for a full charge/discharge [92]. The average operation voltage of 1.7 V of this cell, recorded the second highest for ZIBs through Zn-ions intercalation/de-intercalation mechanism. The same group reported again on the same material of ZnHCF, prepared under different conditions, exhibiting a high energy density of 104  Wh kg−1, providing an operating voltage of

Recent Developments of Zinc-Ion Batteries

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Table 2  Collective summary of V-based electrodes along with their cell configurations and electrochemical performances for aqueous ZIBs Cathode morphologies/ preparative methods V2O5 nanosheets / sol-gel

Electrolyte + additive 3 M ZnSO4

Potential Window 0.4–1.4

Porous V2O5 nanofibers / electrospinning V10O24•12H2O dendrite / hydrothermal Na0.33V2O5 nanowires / hydrothermal Na5V12O32 nanobelts / hydrothermal Na0.76V6O15 nanobelts / hydrothermal K2V8O21 nanobelts / hydrothermal Ag0.4V2O5 nanorods / hydrothermal (NH4)2V10O25•8H2O nanobelts / hydrothermal LiV3O8 flakes / solid-state reaction NaV3O8•1.5H2O nanobelts / liquid–solid stirring method

3 M Zn(CF3SO3)2 3 M Zn(CF3SO3)2 3 M Zn(CF3SO3)2 2 M ZnSO4

0.5–1.5

KV3O8 nanostructure / hydrothermal Zn2V2O7 nanowire / hydrothermal Fe5V12O39(OH)•9H2O nanosheets / water bath Zn3V2O7(OH)2•2H2O nanowire / microwave V5O12•6H2O nanobelts / electrodeposition K0.5V2O5·0.76H2O nanorods / hydrothermal (NH4)2V4O9 sheets / hydrothermal Zn0.3V2O5·1.5H2O nanospheres / in situ electrochemical cycling

Current density (mA g−1) 100

0.7–1.7

588 (1C = 294) 500

0.2–1.6

200

0.4–1.4

500

2 M ZnSO4

0.4–1.4

500

2 M ZnSO4

0.4–1.4

1000

3 M ZnSO4

0.4–1.4

5000

3 M Zn(CF3SO3)2 1 M ZnSO4

0.7–1.7

500

0.6–1.6

133

1 M ZnSO4 + 1 M Na2SO4 1 M ZnSO4/ gelatin 2 M ZnSO4

0.3–1.3

1000

0.3–1.3

500

0.4–1.4

500

1 M ZnSO4

0.4–1.4

300

3 M Zn(CF3SO3)2 1 M ZnSO4

0.4–1.6

1000

0.2–1.8

200

3 M Zn(CF3SO3)2 1 M ZnSO4

0.2–1.6

500

0.4–1.4

8000

3 M Zn(CF3SO3)2 3 M Zn(CF3SO3)2

0.3–1.3

100

0.3–1.6

10,000

Cyclability (mAh g−1) 182 after 30 cycles 166 after 500 cycles ~135 after 500 cycles ~250 after 500 cycles ~200 after 100 cycles ~110 after 100 cycles ~220 after 50 cycles 216 after 2000 cycles ~180 ater 1000 cycles ~140 after 65 cycles 221 after 100 cycles ~120 after 120 cycles ~90 after 1000 cycles 197 after 200 cycles 200 after 50 cycles 101 after 300 cycles 346 after 100 cycles 150 after 1500 cycles 328 after 100 cycles 214 after 20,000 cycles

Ref. [76] [77] [78] [68] [79] [79] [80] [81] [82] [63] [83]

[84] [70] [85] [65] [86] [87] [88] [89]

(continued)

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J. Kim et al.

Table 2 (continued) Cathode morphologies/ preparative methods Mn0.15V2O5·nH2O flower-like / microwave assisted hydrothermal V7O16 nanotubes /

Electrolyte + additive 1 M Zn(ClO4)2 in PC

Potential Window 0.2–1.7

Current density (mA g−1) 10,000

3 M Zn(CF3SO3)2

0.3–1.9

2400

Cyclability (mAh g−1) 153 after 8000 cycles 175 after 950 cycles

Ref. [90]

[91]

1.73  V [93]. Interestingly, the prepared cathode retained 80% discharge capacity in the 3 M ZnSO4 electrolyte at 300 mA g−1 current density over 200 cycles, inferring that the electrode possesses excellent stability. An analogue structure, CuHCF (KCu[Fe(CN)6]), shows a new, safe, and environmentally friendly ZIB with 20 mM ZnSO4 aqueous solution (pH of 6) as the electrolyte through the effects of Zn2+ ions intercalation and H2 evolution [94]. The battery provided ~96% capacity retention over 100 cycles at 1C with an average discharge potential of 1.73 V. The high cycling efficiency is attributed to the deliberate choice of electrolyte pH and the low zinc concentration that ultimately leads to the formation of compact zinc morphology with relatively low surface area at the anode. Upon repeated cycling, the rough surface of the Zn anode facilitates improved Zn-deposition and hence effective long-term cycling by actively preventing the dendritic growth formation and suppressing the corrosion-related hydrogen evolution reaction at the anode. This study proved that by operating the electrode at the appropriate pH, the design of preferentially oriented zinc ZIBs with long-term cycling stability can be realized [95]. Jia et  al. explained the Zn storage mechanism in Cu-HCF electrode in 1 M ZnSO4 electrolyte via a Zn2+ intercalation/de-intercalation mechanism accompanied by solid phase diffusion kinetics (direct interstitial) [96]. In this reaction, the Zn2+-ions intercalate migrate and occupy just the interstitial sites without altering the overall structure of the active material; however, the electrostatic force of the intercalated Zn2+ ions tend to change the binding energies of the ions in its vicinity. In addition, the ions initially occupying their respective interstitial positions present steric hindrance and obstruct the diffusion of the incoming Zn2+ ions, thereby the latter ions require more energy for intercalation. This effect is reflected in the corresponding two discharge plateaus at 0.8–0.5 V and 0.1–0.01 V domains, respectively. Thus, 56 mA h g−1 discharge capacity or nearly 65% of the theoretical capacity is realized. The cell shows discharge capacity of 43 mAh g−1 at 20  mA  g−1 current density over 20  cycles. Further investigations by Kasiri et  al. revealed that the nature of the anion and the Zn-concentration in the electrolyte affects the electrode stability [25]. Among the different electrolytes (ZnSO4, ZnF2, Zn(ClO4)2, and Zn(NO3)2) studied for Cu-HCF, ZnSO4 showed stable cycle ­performance at a given concentration. The (de)intercalation in CuHCF initially follows a single-phase reaction, however, after attaining an equilibrium stage (critical amount of Zn2+ ions in the lattice of the cathode), the phase partially converted into ZnHCF; later insertion follows a two-phase reaction. This phase transition depends

Recent Developments of Zinc-Ion Batteries

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strongly on the current rate; at 10 C the phenomenon is weakly observed only after 800 cycles, while at 1C the transition is already completed after 250 cycles. Thus, understanding the aging mechanism will be helpful to achieve electrode stability under long-­term cycle lifespan.

3.2  Anodes Zn metal is the major anode component used for ZIBs, while other materials like Mo6S8 and Mo2.5 + yVO9 + z were tested as anodes for ZIBs. Chae et al. first demonstrated Zn-ion intercalation into Chevrel phase Mo6S8 using a 0.1 M ZnSO4 electrolyte in a three-electrode beaker-type cell set-up (Fig. 4a) [97]. During the discharge process, Zn-intercalation into Mo6S8 occurred via single-phase and two-phase reactions in the composition range of Mo6S8 to Zn0.24Mo6S8 and Zn0.24Mo6S8 to Zn2.2Mo6S8, respectively (Figs. 4b and c). Within a potential domain of 0.2–1 V, the electrode delivered an initial discharge capacity of 134  mAh  g−1 at 0.05C and retained almost 62% of the initial capacity upon returning from cycling at four progressive current densities from 0.05–1 C, the current rate maintained at each rate for four cycles. Cheng et al. studied Mo6S8 electrode in both aqueous (1 M ZnSO4) and non-aqueous (1 M Zn(ClO4)2 in acetonitrile) electrolytes. The electrodes exhibited good cyclability in aqueous and non-aqueous electrolyte mediums, respectively, as average stable capacities of ~60 and  ~  65  mAh  g−1 was maintained for 150 discharge/charge cycles at 180 mA g−1 [98]. ZnxMo2.5 + yVO9 + z with an open tunnel structure demonstrated good rate capability and cyclability in both aqueous (0.5  M Zn(CH3COO)2 electrolyte) and non-­ aqueous (0.2  M Zn(CF3SO3)2 in 1:4 PC/dimethylsulfoxide (DMSO) solution) electrolytes [99]. The anode exhibited capacities of 180 and 135 mAh g−1 in aqueous and non-aqueous electrolytes, respectively, at 20  mA  g−1. During Zn-ion ­intercalation/de-intercalation, the anode was able to retain the parent structure under electrochemical cycling.

Fig. 4 (a) Illustration of ZnMo6S8 structure. (b) CV curves of the Mo6S8 electrode at 0.05 mV s−1. (c) Initial discharge and charge curves of the Zn-Mo6S8 cell at 0.05 °C (6.4 mA g−1) in the voltage range of 0.25–1.0 V vs. Zn/Zn2+. (reprinted with permission from ref. [97])

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4  Non-aqueous Zinc-Ion Batteries Although the use of aqueous electrolytes, both acidic and alkaline solutions, appears promising, challenges of understanding the reaction mechanism, hydrogen evolution, limited operating potential window, passivation layer formations affecting Zn-ion diffusion and hence cycling ability remain [100]. For example, the role of proton (de) intercalation and Zn-hydroxide precipitate formation (dissolution) in aqueous ZIBs has not yet been completely understood [42]. In addition, the knowledge on the structural variations in the cathode especially with various manganese oxide polymorphs during Zn-ion insertion from aqueous electrolyte medium is limited [36]. Meanwhile, recently, a combination of computational and three electrode based experimental studies confirmed that non-aqueous electrolytes, viz., acetonitrile-­Zn (II) bis(trifluorometylsulfonyl)imide (AN-Zn(TFSI)2), AN-Zn(CF3SO3)2, and propylene carbonate (PC) − Zn(TFSI)2 support reversible Zn-deposition/stripping on Zn metal anode with no irreversible reaction or further redox reaction thus offering anodic stability of electrolyte components of anions and solvents within a wide electrochemical window (~3.8 V vs Zn/Zn2+). Also, the linear sweep voltammetry of the three electrolytes shows highest current density at 0.5 M concentration (Fig. 5a) [101]. More importantly, this breakthrough report facilitates the access of cell potentials beyond the usual aqueous electrolyte window (~ 1.5 V). Moreover, this warrants a critical investigation on non-aqueous ZIBs towards gaining a deeper understanding on their underlying mechanisms and advancements for practical applications. Furthermore, the work on non-aqueous ZIBs with intercalation cathodes are only in their infant stages and more strides forward are expected to be taken in this direction soon. As reported for aqueous ZIBs, the strategies of utilizing layered-type compounds with apparently low migration energies (than 3-D spinel or 1D-olivine structures) and designing the corresponding nanostructures incorporated with “pillar”

Fig. 5 (a) Linear sweep voltammetry curves at a scan rate of 0.025 V s−1 for 0.5 M AN−Zn(TFSI)2, AN−Zn(CF3SO3)2, and PC − Zn(TFSI)2 electrolytes to evaluate the stability of anions and solvents with respect to Zn anode through a three electrode based analysis. (reprinted with permission from ref. [101]) (b) Schematic illustration of a non-aqueous ZIB system, i.e., Zn||0.5  M AN− Zn(TFSI)2||V2O5 (red: oxygen (O), green: zinc (Zn), blue: vanadium (V)) showing intercalation/ de-intercalation of Zn2+ ions into the bilayered vanadium oxide. (c) Electrochemical profiles of the non-aqueous ZIB within the potential window of 0.3-1.5 V at 0.1 °C. (reprinted with permission from ref. [26])

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elements/molecules enabling multivalent ion migration within the crystalline hosts are promising. To this end, Senguttuvan et al. developed a 0.85 V Zn-full cell based on Zn metal anode and a hydrated bilayered V2O5 electrodeposited on a carbon substrate as the cathode and a solution of 0.5 M Zn(TFSI)2 in acetonitrile as the non-­ aqueous electrolyte. The schematic illustration of the non-aqueous ZIB is provided in Fig. 5b. Electrochemical profiles obtained for the 0.3-1.5 V window are provided in Fig.  5c. The cycle performance studies revealed stable specific capacities (~170 mAh g−1 for over 120 cycles at 0.1 °C) with 99% Coulombic efficiency and suggested the high structrual integrity of the cathode thereby demonstrating the excellent plating/stripping effectiveness in the corresponding electrolyte. More importantly, a 130  mAh  g−1 specific capacity at 20  °C rates corresponding to 1500 W kg−1 power that is almost equivalent to standard Li-ion batteries was realized. The enhanced electrochemical performance was attributed to the layered cathode structure being stabilized by water molecules at the interlayers (with wide spacing of 11–13  Å) thereby enabling facile reversible Zn-ion (de) insertion. Importantly, the oxide host structure remained unaffected by the presence of oxygen-bearing TFSI salt in the electrolyte [26]. Han et  al. studying the insertion chemistry in a layered-type water stabilized δ-MnO2 nanostructured cathode (K0.11MnO2·0.7H2O) confirmed reversible Zn insertion in a non-aqueous electrolyte media of Zn(TFSI)2 in acetonitrile solution without macroscopic phase transitions and significant proton intercalation. The fabricated Zn-δMnO2 cell delivered a maximum of 123 mAh g−1 over 125 cycles at 0.1 °C with gradual capacity fade in the potential domain of 0.05–1.9 V. The capacity fade was attributed to the reduced electrolyte performance, the competing reactions for ZnO formation on the electrode surface and Mn dissolution [102]. Cubic PBAs with formula, AxM1[M2(CN)6]y·nH2O, (M1/M2—metal ions; cyanide (CN) ligands; A—mobile alkaline metal ions) belonging to Fm3̅m space group benefit as cheap, nontoxic, easy-to-prepare, and open-framework cathodes of ZIBs. An eco-friendly PBA cathode, K0.05Fe(III)[Fe(III) (CN)6]·2.6H2O, integrated with a cheap, biodegradable, and biocompatible “ionic liquid in water” electrolyte and a nontoxic/low-cost Zn anode delivered reversible discharge capacities of 120 mAh g−1 at 0.1  °C with 99% Columbic efficiency for the initial ten cycles at a current of 10 mA g−1 (∼0.1 °C) in the potential domain 0.8–2 V. Fourier electron density analysis with powder XRD studies done on another PBA cathode, potassium nickel hexacyanoferrate K0.86Ni[Fe(CN)6]0.954(H2O)0.766 in 0.5 M Zn(ClO4)2 in acetonitrile electrolyte medium confirmed that the inserted Zn2+-ion occupies the center of the large interstitial cavity, the same location as of the potassium and zeolite water in the cubic open-framework structure [103, 104]. Non-aqueous gel polymer electrolytes (GPE) based on ionic liquids, 1-buthyl-­3methylimidazolium trifluoromethanesulfonate (BMIM triflate), and 1-ethyl-3-methylimidazolium bis(trifluoromethylsulfonyl) imide (EMIM TFSI), respectively, were prepared for use in ZIB applications. The inclusion of the ionic liquid led to the formation of good GPEs while the retention of NMP molecules in the polymer was crucial to improve GPE properties. In the Zn-MnO2 cell with the former GPE (PVdF-HFP/BMIM triflate/Zn(CF3SO3)2), ex situ XPS, SEM, and EDX studies

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c­ onfirmed the initial co-insertion/de-insertion of Zn-ion and triflate anions during discharge/charge reaction. A discharge capacity of 120 mAh g−1 was achieved by the former Zn//IL-GPE//MnO2 battery at 6.25 mA g−1. These studies are still at the infant stages and it is obvious that further research is currently underway to fully exploit the usefulness this system [105, 106].

5  Summary and Outlook Rechargeable ZIBs utilizing both aqueous and non-aqueous electrolytes have advanced quite significantly ever since its renewed interest from 2012. In specific, the aqueous ZIBs are highly attractive due to the benefits of their high safety, low-­ cost, and less toxic features. As presented above, the identification of high-­ performance electrodes, especially cathode materials, have been crucial for battery performance enhancements to meet practical stationary storage applications. Since the major find for the aqueous cathodes have been focused on manganese and vanadium-­based materials, the Ragone plot for few of the important materials studied so far have provided in Fig. 6a, b, respectively, to be used for their evaluation for practical suitability. Among manganese-based electrodes, different polymorphs of nanostructured MnO2 and their composites (formed from electrically conducting inclusions) prepared by a variety of syntheses were developed to realize improved performance. For example, among the manganese-based electrodes for ZIBs ­application, β-MnO2 nanorod delivers a high specific energy of 254/110 Wh kg−1 at a specific power of 197/5910 W kg−1 [38]. In the family of vanadium-based electrodes, Zn0.25V2O5 electrode exhibits specific energies of ~250 and ~ 150 Wh kg−1 (based on cathode mass) at the specific powers of ~40 and 4860 W kg−1, respecb -1

Specific energy (Wh Kg )

Specific energy (Wh Kg-1)

a 100

10

1

Na2V6O16·3H2O LiV3O8 V2O5 VS2 Na3V2(PO4)3 100

V2O5·nH2O/rGO Zn0.25V2O5·nH2O Zn2V2O7 H2V3O8 Ca0.25V2O5·nH2O

1000 Specific power (W Kg-1)

10000

100

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β−MnO2 α−MnO2 δ−MnO2 λ−MnO2 ε−MnO2 100

Mn2O3 Mn3O4 Todorokite MnO2 Spinel ZnMn2O4 Ramsdellite MnO2 1000 Specific power (W Kg-1)

10000

Fig. 6  Ragone plots for few selected for (a) vanadate based cathodes for aqueous ZIBs application including Na2V6O16.3H2O [107], LiV3O8 [63], V2O5 [108]. VS2 [74]. Na3V2(PO4)2 [75]. V2O5. nH2O/rGO [67]. Zn0.25V2O5..nH2O [62]. Zn2V2O7 [70]. H2V3O8 [66]. and Ca0.25V2O5.nH2O [71]. (b) Manganese-based cathodes for aqueous ZIBs application including β-MnO2 [38], α-MnO2 [50], δ-MnO2 [37], λ-MnO2 [40], ε-MnO2 [42], Todorokite MnO2 [46], Ramsdellite MnO2 [109]. Mn2O3 [47]. Mn3O4 [48]. and Spinel ZnMn2O4 [41].

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tively [62], while Zn2V2O7 nanowire displays the specific energies of 166.3 and 114.3 Wh kg−1 at specific powers of 36 and 3168 W kg−1 (based on cathode mass), respectively [70]. This nanowire Zn-ion battery offering an energy density of 87.3/60 Wh kg−1 for a given power density of 18.9/1663 W kg−1 at 50/4400 mA g−1 current density (based on whole cell) is still higher than those of commercial Pb-acid (~30 Wh kg−1) and Ni-Cd (50 Wh kg−1) batteries. Thus, the performances of the electrodes, mainly based on manganese and vanadium materials discussed above, show potential for clear superiority to the performance parameters of commercially used alkaline Zn/MnO2 batteries, the scope for further enhancements on ZIB performance is still wide open. The following are the conclusions that can be derived from the research developments on ZIBs so far: 1. Tunnel-type and layered-type electrode structures with sufficiently wide geometric dimensions tend to favor reversible Zn insertion; though some of these electrodes undergo irreversible phase transition that seem to be dependent on the crystal structure, synthesis, and morphology of the electrode and the type of electrolyte used. 2. The strategies of conductive inclusions combined with pre-inclusion of additives in nanostructured manganese-based electrodes and structural water inclusion and/or metal-ion doping in vanadium-based electrodes have been successful, at least to a certain extent, to avoid active material dissolution, and hence increase zinc storage capacities in addition to extending the cycling stability of these electrodes to a few thousands of cycles from a few hundred cycles under high current drains. 3. Although major strides on the electrochemical performance of ZIBs have been made, three major drawbacks including low intrinsic electrical conductivity of manganese/vanadium materials, active material dissolution, especially, in manganese-­based electrodes and formation of by-products via electrode phase transition or side reactions hinder the achievement of practical zinc storage capacities equivalent to theoretical values. 4. Compared to other salts, ZnSO4 and bulky anion salts including Zn(CF3SO3)2 and Zn (TFSI)2 show high solubility in aqueous solutions and form mildly acidic electrolytes that demonstrate good electrochemical stability, minimum dendritic growth and cause considerably low corrosion in ZIBs. ZnSO4 is low-cost; however, the formation of byproducts like ZHS during can accumulate on the cathode surface thereby leading to capacity fade during repeated cycling. Whereas the use of Zn(CF3SO3)2/Zn (TFSI)2 salts offers higher electrochemical stability and performance but at the cost of the economic feature as the salt price is highly expensive than the low-cost ZnSO4 to realize large-scale production. For non-­ aqueous ZIBs, only Zn (TFSI)2 salt has been used so far to demonstrate reversible Zn insertion. 5. As for anodes, more efforts to analyze the pH of the solution can be a useful approach to control the hydrogen evolution and avoid the formation of dendrites thereby offering stability to repeated Zn stripping/plating reactions for the long term.

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Therefore, in less than a decade, the research and development on ZIBs has been very fruitful. From the abovementioned discussion the perspectives on the development of state-of-the-art ZIB cathodes can be summarized as follows: 1. Although high performances in terms of high zinc storage capacities and long-­ term cycling stabilities under high applied current densities are achieved for ZIBs, there are challenges in identifying the underlying the electrochemical reaction and the by-products. Also, it is worth noting that the monitoring of electrochemical reactions through low current drains under long-term cycling is essential for such aqueous systems to be realized for practical stationary applications. At present, the electrochemical reactions proposed are contrasting and highly debated. For example, in MnO2, the zinc storage mechanism was described by Zn-intercalation/de-intercalation without or with additional phases and a conversion reaction, respectively. Even within the short time duration of writing and publishing this chapter, additional reaction mechanisms based on reversible co-­ intercalation reaction of H+/Zn2+, either sequentially or concordantly, combined intercalation-conversion reactions have been proposed [54]. Further, the origin and the role of discharged by-products on the electrochemical reaction remains evasive. Hence, a balanced analysis of different real-time spectroscopic measurements based on electrochemical, X-ray, vibrational and NMR techniques in combination with computational techniques can help to confirm the reaction mechanism in ZIBs. 2. Special focus on the understanding of electrode degradation mechanism and interface reactions related to the solid-electrolyte-interface layer in detail via in situ electrochemical techniques including potentiostatic electrochemical impedance spectroscopy (PEIS) or galvanostatic intermittent titration technique (GITT) during on-time cycling will be required to establish their influence on the electrochemical reaction and shed more light on the aging issue of the ZIB during electrochemical reaction. Also, these techniques will help in answering many questions in understanding the exact role of structural water in layered vanadium cathodes. 3. Focus on identifying nanostructured materials based on phosphates or amorphous or organic/inorganic hybrid materials will also be the norm to achieve further performance improvement in the ZIB battery system. The study on phosphates is motivated by their unique 3D open-framework structures with transition-­metal ions and covalently bonded polyanion (PO4)3− units. In addition, they demonstrate high structural stability, tunable operating potential window via altering local environment, high thermal stability and can accommodate oxidation reactions at high charge potentials [110]. Specially, the use of fluorophosphates can benefit in raising the operating potential window in aqueous ZIBs. On the other hand, amorphous materials can evade macroscopic phase transitions during electrochemical reaction and thereby preserve the electrode from abuse related to volumetric changes and facilitate long-term electrode cycling stability.

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4. The study of more chalcogenide electrodes with less electronegative elements can be advantageous than highly electronegative oxygen-containing transition-­ metal oxides and improve the sluggish Zn2+-ion diffusion kinetics and enhance the performance parameters of ZIBs. Different strategies to developing specially designed or hierarchical structured nanocomposites will be made to effectively suppress cathode dissolution. Specifically, in relation to non-aqueous ZIBs, being only just 1 year since its introduction, there will be a stiff race towards achieving complete advancement. 5. The feasibility of widening the operating potential window (> 1.5 V) for ZIBs can be realized by electrolyte optimization. For example, the effect of various electrolyte parameters including salts, additives and their concentrations, pH and their influence on the electrochemical reaction needs to be unraveled via on-time monitoring of these factors. 6. One of the immediate challenges that should be addressed for the anodes is the inhibition of dendrite formation. The dendrite factor becomes significant when fabricating large-scale ZIBs beyond lab-scale limitations. This could be achieved by suppressing the hydrogen evolution reaction during electrochemical reaction. Surface coating of polymers or protective layers like an alternative SEI layer in situ formation on the anode by using appropriate additive inclusions in the electrolyte can be helpful to prevent dendrite formation [111]. Another method to evade dendrites is the use of solid/gel electrolytes with high mechanical and ductile properties. This can also contribute to high zinc storage capacities and performances [112]. 7. From the long-term economic viewpoint, it is crucial for a new emerging energy economy to be centered on cheap, green, and sustainable battery systems. The low-cost and environmentally friendly features make manganese-based electrodes attractive for use in ZIBs. But the case is not true for vanadium as it is faced with environmental issues. In economic terms, the recent years have seen a significant fluctuation in the raw material costs. Specially, the cheapest raw material of vanadium, V2O5 was priced the lowest in 2016 (~ USD2/lb) before shooting up to almost 17 times by late 2018 (~ 33.2 USD/lb). Currently, vanadium price is trending at 6.32 USD/lb., which is still three times higher than the lowest cost. This uncertainty has become a matter of concern for the use of vanadium-based materials, in general, for rechargeable batteries including ZIBs. Therefore, it is immediately essential to develop efficiently reliable and eco-­ friendly methods for large-scale economical production of V2O5 [113]. 8. As the world of portable electronics is upgrading to next-generation technologies using wearable and flexible rollup displays and power devices, there will be a significant effort to develop wearable or flexible aqueous ZIBs. From the engineering point of view, many more new strategies to design the ZIB in various shapes and types will be pursued. More importantly, major stress will be given to incorporate environmentally safe components and follow low-cost approaches to realize stable, safe and economical grid-scale energy storage applications.

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Overall, it is obvious that research on ZIB technology, aided by good engineering, can demonstrate better performance, economic and environmental advantages than the present-day Pb-acid and Ni-Cd batteries. The ZIBs clearly hold the benefit to serve the demands for green and sustainable energy storage in modern society. Acknowledgements  This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (MIST) No.2020R1A2C3012415 and 2018R1A5A 1025224.

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Molten Sodium Batteries Erik D. Spoerke, Martha M. Gross, Stephen J. Percival, and Leo J. Small

Abstract  The rapid growth in demand for electrical energy storage in the twenty-­ first century has renewed interest in molten sodium batteries for safe, reliable, typically large format storage. These batteries take advantage of globally abundant sodium as a key active material in the system, employing solid-state ceramic separators, and utilizing several different cathode chemistries in battery design. This chapter introduces these battery systems, primarily focused on sodium-sulfur and sodium-nickel chloride (ZEBRA) batteries. It describes the materials chemistries of the anodes, the separators, and the cathodes of these batteries, highlighting both virtues and challenges to be overcome with further research and development. It further highlights new chemistries and materials that may enable key innovations to battery design that will lower the operational temperatures, improve safety, and drive down limiting costs of these systems. Ultimately, improvements resulting from these research and development efforts are expected to expand an already growing global market for large-scale molten sodium batteries. Keywords  Large format batteries · Molten sodium · High temperature · Sodium-­ sulfur · Sodium metal halide · ZEBRA · Sodium-air · Molten salt · Solid-state electrolyte · β”-alumina · NaSICON · Ion conducting glasses · Borates · Ionic liquids · Sodium battery deployment

1  Introduction 1.1  Brief History Metallic sodium (Na) batteries, utilizing a molten sodium anode, have been an active area of research and development since the 1960s. In 1968, the sodium-sulfur (NaS) battery was patented by Ford Motor company, who was pursuing it as a candidate for automotive applications [1]. The sodium metal halide battery, known E. D. Spoerke (*) · M. M. Gross · S. J. Percival · L. J. Small Sandia National Laboratories, Albuquerque, NM, USA e-mail: [email protected] © National Technology & Engineering Solutions of Sandia, LLC 2021 M. Alston, T. N. Lambert (eds.), Energy-Sustainable Advanced Materials, https://doi.org/10.1007/978-3-030-57492-5_3

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more commonly as the ZEBRA battery (Zeolite Battery Research Africa Project or more recently, Zero-Emission Battery Research Activities), was originally patented in 1975 and was pursued as an alternative to the NaS battery [2]. Until recently, however, these technologies have found relatively little utility, despite tremendous potential technical value. Historical problems with battery cost, safety, operational temperature, and long-term performance have resulted in relatively limited demand for these kinds of batteries. Over the past 10–20  years, however, the explosive increase in demand for the storage and distribution of electrical energy has motivated a renewed interest in a variety of battery technologies. There is clear opportunity for sodium-based batteries to fill critical gaps in the current electrical energy storage technology portfolio. Generally speaking, there are many different sodiumbased battery technologies emerging to meet the varied global demand for batteries, ranging from sodium-ion and flow batteries to solid-state systems. This chapter, however, will specifically focus on molten sodium batteries with an emphasis on the materials science of these promising battery systems. It will introduce existing battery chemistries, highlight some of the challenges with the current state of the art, and discuss opportunities to advance these batteries in the coming years.

1.2  Battery Development Considerations Despite the focus of this chapter on molten sodium batteries, it is important to note that the integration of batteries into the emerging global energy storage future will almost certainly include a multitude of chemistries and technologies as no single battery technology is “right” for every application. Batteries are essentially electrochemical reactors, and both the advantages and limitations of each chemically distinct reactor should be considered when selecting a battery system for a particular application. Requirements related to how much energy a battery can store and how quickly, or slowly, the system can be charged and discharged matter a great deal. The functional lifetime of the battery, particularly for secondary (rechargeable) batteries, is also important, affecting not only performance, but also the effective cost of a system. Factors such as a battery’s size or weight, if it will be stationary or mobile, and the climate where the battery will be used should also be taken into consideration. The safety of the battery is an important factor as well and has drawn substantial interest recently. It is important to ask if there are any hazardous side reactions to be aware of, either during normal operations or in the event of an unexpected assault on or failure within the battery. Widely publicized fires from failing lithium-ion batteries, for example, highlight the importance of this issue, both for applications in ubiquitous personal electronics and for very large grid-scale systems with much higher consequences of failure. It should be noted that although molten sodium batteries were originally conceived and developed for transportation applications, current battery development of these technologies is focused largely on grid-scale energy storage. Current research has an eye on all of these factors, with different chemistries providing or emphasizing different solutions based on application-­dependent requirements.

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All of the issues above are ultimately tied to both battery performance and the critically important cost of a battery, weighed against the market demand for the battery. An extremely high performing sodium battery composed of overly costly materials or that is too expensive for utility companies or industrial interests to deploy is likely to have limited impact on the need for energy storage. In contrast, an inexpensive battery that has a short lifetime, poor performance, or is unsafe, is ultimately not cost-effective either. In the discussions below, it is clear that researchers are looking not only for materials and chemistries that make energy storage possible, but there is clear consideration of how the materials chemistry in each system make these batteries practical. One of the most important factors limiting the widespread practical application of molten sodium batteries is the reality that these batteries are typically utilized at elevated temperatures, near 300 °C. The high operating temperature keeps the battery above the melting temperature of the necessary molten components in the system, but it also facilitates rapid ion transport and reaction kinetics that are important to delivering reasonable power from the storage device. This high temperature, though, can lead to accelerated aging or degradation of material components, leads to potentially detrimental (or dangerous) side reactions, requires more advanced (often expensive) thermal management strategies, limits the application space of the system, and overall drastically increases the cost of the battery, even when the basic cell components might be inexpensive. One of the primary challenges to modern battery researchers is to lower the operating temperature of the battery without sacrificing the performance traditionally seen at elevated temperatures. Ultimately, the goal remains to develop a cost-effective molten sodium battery system suitable for widespread grid-scale application.

1.3  Battery Basics To understand sodium batteries it is important to first understand how a battery is put together. In its simplest form, schematically illustrated in Fig. 1, a battery consists of an anode (negative electrode) and a cathode (positive electrode), positioned opposite one-another and separated by a physical and electronic barrier (polymer or Fig. 1 Schematic depiction of a molten sodium battery comprising a molten sodium anode, a solid-state sodium-ion separator, a molten (or partially molten) catholyte, and a cathode current collector. (Sodium serves as anode current collector)

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ceramic membrane) and ion-conductive electrolytes (ceramic or other solid-state electrolytes often also act as the separator). Current collectors, interfaced with a battery’s anode and cathode, facilitate electron transfer in and out of the battery. In a molten sodium battery, the anode is sodium metal, and the battery must be operated above its melting temperature (97.8  °C). This molten sodium anode wets a sodium-ion conducting ceramic separator that isolates the anode physically and electronically from various cathode chemistries that will be discussed in greater detail later in this chapter. As the battery discharges, an oxidation reaction at the anode extracts electrons from the metallic sodium, creating Na+ ion. The extracted electrons are collected by the current collectors and provide power as they are shuttled through an electrical circuit to the current collector at the cathode where they participate in electrochemical reductions, dependent on the specific cathode chemistry. For example, in a sodium-sulfur battery, molten sulfur is reduced to form molten polysulfides, while in a ZEBRA battery, Ni2+ ions are reduced to metallic nickel. Meanwhile, the oxidized sodium ions (Na+) must cross through the ion conducting separator and/or electrolytes to the cathodic side of the battery, where they participate in charge-balancing electrochemical reactions. Together the external transport of the electron from anode to cathode and the internal transport of the oxidized ion complete the electrochemical circuit of the discharging battery. In the case of rechargeable (also known as secondary) batteries, this process is reversed to recharge the battery. One complete charge/discharge process is known as a “cycle.” The energy density of a battery is the amount of energy in a given system, per unit mass (gravimetric energy density, also called specific energy) or per unit volume (volumetric energy density). This energy density is the product of its capacity (how much charge the battery can effectively charge and discharge) and its voltage. The maximum ideal voltage the battery is capable of achieving is determined by the free energies of the electrochemical reactions at the anode and cathode. The voltage, and therefore the energy density of the battery, is strongly dependent on the chemistry within the battery. The power of a battery is, effectively, how quickly the battery can release its energy. This power can also be normalized by weight or size to determine the power density of the system. It is worth noting that a “high-power” system may not necessarily have a high energy density; a small amount of energy delivered very rapidly would be considered a high-power system. One of the common goals of battery researchers in general is to create a system that combines high-power and high energy density, but not all battery applications require these capabilities. Figure 2 relates different grid-scale electrical energy storage technologies’ power ratings and discharge times, highlighting the application space where these technologies may provide meaningful utility. On this plot, molten sodium batteries are represented by the NaS and the Na-NiCl2 batteries, which provide sufficient energy storage to discharge on the order of hours (often 4–6) with reasonable power capabilities for applications such grid support and load shifting, but they are clearly not ideal for all applications. It should be noted that the application categories are not absolutely defined by the boundaries shown; significant overlap in these applications is common and dependent on the specific needs of a system

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Fig. 2  Generalized comparison of discharge time and power rating for different electrical energy storage technologies. Data obtained from Ref. [3] and Ref. [4]

or storage community. Still, the combination of long cycle life, low-cost materials, and significant energy density, means that these molten Na batteries have significant potential for select grid-scale energy storage applications. Battery manufacturers, such as NGK and FZSoNick (FIAMM), have deployed some NaS and NaNiCl2 grid-scale batteries around the world for renewables integration, load shifting, frequency regulation, backup power, and other grid-scale applications [3, 5, 6]. Continued expansion of these technologies will be facilitated by significant reductions in capital costs, reductions likely made possible through innovations in materials and cell designs.

2  Battery Components 2.1  Sodium Anode With a redox potential of ENa + / Na  = −2.71 V versus a standard hydrogen electrode, a metallic sodium anode can be coupled with a number of cathode chemistries to achieve useful energy densities, and the global abundance and availability of sodium, relative to lithium, for example, makes it an economically and geopolitically desirable choice [7]. As a metal, it is inherently electrically conductive, making it relatively easy to inject or extract electrons from the anode during electrochemical cycling. Finally, in the molten state, the sodium will naturally flow

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and conform to variable geometries required in battery design. Importantly, this means that during electrochemical discharge, gaps at the sodium-separator interface (a potential challenge for solid-state metal batteries) can be avoided. Batteries that employ solid-state sodium anodes represent an area of growing interest, though the relatively high reactivity of sodium against organic catholytes has made solid sodium anodes impractical for consideration in many sodium-ion systems, and the comparatively low energy density (relative to solid state Li, for example), has limited research efforts into solid-state sodium batteries. Future efforts may see these technologies realize greater opportunities with time. The present work is focused on molten sodium batteries, which are more mature technologies today. The molten metal anode does, however, introduce a number of challenges. First, the batteries must be operated at temperatures above the 97.8 °C melting temperature of sodium. Historically, molten sodium batteries have been operated closer to 300  °C to accommodate other material requirements of the batteries, which has made the molten character of the battery a minor concern. A more common concern is that metallic sodium is strongly reactive with water, which requires that the anode must be handled in a dry or inert atmosphere until the battery is hermetically sealed. By assembling batteries in a discharged state where the sodium is tied up in the cathode chemistry, it is possible to significantly reduce (or potentially eliminate) the amount of sodium metal that must be handled during battery assembly [8]. As a liquid anode, there can also be concerns with wetting of the anode at the anode-­ separator interface. This has been a well-documented problem with traditional separator materials, such as β”-alumina solid electrolyte (BASE), though a number of reports have described methods to address this issue through chemical modification of either the BASE surface or the metallic anode composition [9–15]. As this issue of sodium wetting is not unique to BASE [16], it must be considered for sodium battery development for a variety of systems, particularly for lower temperature applications.

2.2  Separators The sodium-ion conducting membrane that separates the anode and cathode remains one of the most important aspects of sodium batteries and presents some of the greatest opportunities to advance molten sodium batteries. Not only must this serve as a robust electrical and physical barrier between the two electrodes, it must also provide high ionic conductivity while maintaining good chemical and thermal stability in contact with both molten sodium and various, often chemically aggressive cathodic chemistries. Here we give an overview of common separators that have been traditionally used in molten sodium batteries and brief discussions of promising new separator materials.

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BASE Original molten sodium batteries utilized a β”-alumina solid electrolyte (BASE) separator, and many sodium batteries still utilize this material today. BASE has been discussed extensively in the literature if details beyond the scope of this discussion are of interest to the reader [12, 17–20]. BASE may be thought of as Na2O-doped Al2O3 in a rhombohedral (R3m) crystal structure, composed of alternating layers of densely packed Al2O3 and loosely packed Na2O layers, as can been seen in Fig. 3. Sodium ions are able to move very rapidly along loosely packed layers, in what are referred to as conduction planes (or conduction slabs) [20]. Figure 3 also highlights the differences in structure between β-Al2O3 and β”-Al2O3, where subtle changes in the planar stacking sequence and the density of sodium carriers in the conduction planes allows the β” phase to be much more conductive. These properties make BASE an effective ion conductor and separator, particularly at elevated temperatures. At 300 °C, a temperature range in which many molten sodium batteries are operated, the ionic conductivity of polycrystalline BASE is 2–4 × 10−1 S cm−1 [12, 17]. The materials are typically made from inexpensive starting materials and can be manufactured using a variety of different techniques, including traditional solid-­ state chemistry, sol-gel, co-precipitation, freeze-drying, flame pyrolysis, microwave synthesis, and mechanochemical methods [17, 19]. The synthesized materials can then be shaped into the desired form using isostatic pressing, electrophoretic deposition, slip-casting, or extrusion. These materials must then be fired at relatively high temperatures (e.g., ≥1600 °C) to achieve suitable density, mechanical strength, and ionic conductivity. Firing at these high temperatures, however, introduces challenges in controlling ceramic microstructure such as grain size and crystal chemistry, and composition due to

Fig. 3  Projection of the (a) β-alumina and (b) β”-alumina unit cells on (11–20) showing stacking sequence and conduction planes. Reprinted from Ref. [12], Copyright 2010, with permission from Elsevier

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sodium volatility. Poor control over the microstructure can significantly affect the mechanical properties of the BASE, which is important for several reasons. First, the BASE is required not only to separate the anode and cathode electrically, but it must also serve as a robust physical barrier between the two sides of the battery. Allowing physical contact between the highly reactive molten sodium and the catholyte leads to failure of the cell and, for example, with NaS batteries, can lead to violent and hazardous reactions. To ensure the mechanical integrity of these materials they must be made relatively thick and quite uniform. The uniformity represents a processing challenge, while the thickness impacts the overall conductance of the membrane. Introducing resistive elements, such as thick BASE membranes decreases overall cell performance and requires undesirable operation at relatively higher temperatures (≥300  °C) where membrane conductivity is sufficiently high. The introduction of zirconia into the BASE has shown promise in increasing the mechanical properties of the membranes, potentially allowing for higher performing batteries or lower temperature operations, though the introduction of this secondary phase may impact ionic conductivity, particularly at higher concentrations [17, 19, 21]. NaSICON NaSICON (Na Super Ion CONductor) ceramics have been considered as attractive candidates for Na-based batteries, particularly for lower temperature applications. Although some reports have described the use of BASE for reduced temperature batteries, the conductivity of NaSICON below 150 °C is generally considered to be higher than that of traditional BASE separators [10, 11], (Fig. 4). This difference can vary as a function of the quality and composition of each ceramic. The typical sodium-conducting NaSICON is described as Na1+xZr2P3-xSixO12 (0 ≤ x ≤ 3), and forms a rigid hexagonal framework of zirconia (ZrO6) octahedra corner-linked to silica and phosphate ((Si,P)O4) tetrahedra, and contain the so-called

4 Ln ( s T ) / S cm-1 K-1

Fig. 4  Comparison of the ionic conductivities of NaSICON and BASE. Arrhenius plot of Na+ conductivity vs. inverse temperature for both. For reference, room temperature is near a value of 3.4 and 150 °C is near a value of 2.1 on the x-axis. Reprinted from Ref. [22], Copyright 2017, with permission from Elsevier

2 0 -2 -4 -6

NaSICON b"-Al2O3 2

2.5 3 1000/T / K-1

3.5

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M1 and M2 interstitial positions filled with sodium ions (Fig. 5) [20, 23, 24]. The conduction of alkali cations through the NaSICON crystal structure is believed to be based on the movement of ions from site to site through the crystal lattice via channels (bottlenecks) constricted by zirconia octahedra and phosphate or silica tetrahedra. Compositions range from the “P3” composition (x = 0) to the “Si3” composition (x  =  3), though the highest selective conductivity compositions are generally reported near x = 2 [20]. Substitutions on the NaSICON lattice can be used to vary the charge distribution, defect chemistry, and critically, the size of channels and bottlenecks in the crystal lattice, and this approach has been used to improve the conductivity of select NaSICON compounds [25]. NaSICON can be prepared using a number of approaches, ranging from traditional solid-phase ceramic syntheses to sol-gel and microwave syntheses [26]. Among the challenges to synthesizing quality NaSICON is controlling the formation of secondary phases, including zirconia, silicate, and glassy inclusions as well as a poorly defined grain boundary phase [22]. These secondary phases can impact the ionic conductivity as well as the chemical stability and mechanical properties of the solid electrolyte. Highly dense NaSICON with proper composition and minimal secondary phases present, however, can be expected to be highly conductive with values >1 × 10−3 S cm−1 at room temperature and stable against molten sodium. Although there has been some historical debate about the stability of NaSICON exposed to sodium at high temperatures (>300  °C) [27–31], recent reports have shown that NaSICON appears to be “quite stable” at moderate temperatures [16, 22].

Fig. 5  Crystal structure of NaSICON showing M1 and M2 interstitial positions. Reprinted from Ref. [23], Copyright 1997, with permission from Elsevier

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Glasses Glass separators have been studied for molten sodium batteries almost since the development of the sodium-sulfur battery [32]. Glasses offer advantages over ceramic electrolytes in that they exhibit isotropic properties, they do not have grain boundaries, and they are easy and inexpensive to produce in a variety of different form factors [33]. Early work looked primarily at oxide glasses, such as sodium borate glasses and NaSICON derivative NASIGLAS [33, 34]. However, research in this area stalled due to the low conductivities of the glass materials produced (~1 × 10−3 S cm−1 at 300 °C), as well as a lower corrosion resistance compared to BASE [33]. Interest in glass electrolytes has been revived in part due to desire on the part of researchers to improve the room temperature conductivity of the solid electrolyte, with an eye on both all-solid-state batteries as well as low and intermediate temperature molten sodium batteries. Recently, new glass electrolytes have been demonstrated with nominal composition of A2.99M0.005OCl1-x(OH)x where A = Li or Na and M = Ba or Ca [35, 36]. These glasses have high conductivities at room temperature which were demonstrated as being >1 × 10−2 S cm−1 [35, 36]. These new types of ion conducting glasses have been successfully incorporated into lab scale solid-state batteries [36], but have yet to be shown stable against molten sodium metal. Sulfide glasses have shown promise of high ionic conductivities at room temperature compared to oxide glasses, owing to the weaker bonding between sodium and sulfur ions than oxygen in the glass [37]. One disadvantage of sulfide-­ based glasses is the need to handle them in dry atmosphere owing to their highly hygroscopic, reactive nature [38]. Glasses can be produced both by melt quenching and by mechanical alloying. In mechanical alloying, precursor materials are mechanically mixed until molecular bonding of the materials occurs by a mechanochemical reaction to form a powder. The powder is subsequently pressed to the desired form [37, 38]. The predominant sulfide glasses studied for use in batteries consist of Na2S–GeS2, Na2S–P2S5, and mixed Na2S–GeS2–P2S5 which have room temperature conductivities ranging from ~1 × 10−5 to 1 × 10−6 S cm−1 depending on the exact composition of the glass [39]. Glasses containing GeS2 typically are composed of GeS4 tetrahedral units, while those containing P2S5 are thought to consist of a mixture of PS4 tetrahedra, P2S7 with bridging sulfur, and P2S6 with P–P bonds [38, 40, 41]. The mechanical and electronic properties of the glasses are highly dependent on both composition as well as the synthesis and forming methods [38–40]. Research in this field has so far largely focused on material conductivity and mechanical properties, as well as application to room temperature all-solid-state batteries. More work is needed to determine the long-term chemical and electrochemical stability of these glasses with molten sodium and many of the cathode chemistries discussed below across a range of operating temperatures.

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Glass Ceramics The conductivities of sulfide glasses exceed that of their oxide counterparts, but remain low compared to ceramic electrolytes. This drawback cannot fully overcome their superiority in terms of manufacturability and, for this reason, more recent work has explored glass-ceramic composites which combine some of the ease of glass production with the increased conductivity expected from a ceramic. To this end, work has largely focused on the development of the cubic Na3PS4 phase within a glass matrix owing to its high ionic conductivity of ~2 × 10−4 S cm−1 at room temperature [42]. Glass ceramics are typically produced by the heat treatment of glasses, such as the ones discussed previously, to precipitate a ceramic phase. Cubic Na3PS4 is generally considered to be a high temperature phase, while its lower conductivity tetragonal phase is considered to be the low-temperature phase. However, it has been found that the applications of moderate temperatures can precipitate cubic Na3PS4 within a glassy matrix with minimal formation of the tetragonal phase. It has been shown that care must be taken, as treatment at higher temperatures can still result in the precipitation of the low-temperature tetragonal Na3PS4 [42–44]. Study of the Na3PS4 ceramic alone has indicated that halogen doping may help to further improve the conductivity of the Na3PS4 [45]. Other work has also shown that substitution of the S for Se, and substitution of Sb or As for P can also raise the conductivity of the ceramic phase [5, 46]. Work on lithium-based analogues has indicated good stability with Li metal, but further work is necessary to determine the long term stability of the sodium glass ceramics with molten sodium anodes [19]. Polymers Polymer-based separators are another area of interest due to the low cost of materials and ease of manufacturing. The use of polymer separators for molten sodium batteries is almost completely unexplored as traditional battery chemistries require operating temperatures in excess of 300 °C, which exceeds the functional operation temperature of most polymers. However, with the development of new cathode chemistries with lower operating temperatures, polymer separators are becoming attractive candidates for these batteries. Poly(ethylene oxide) (PEO) has been studied as a Na+-conducting polymer electrolyte in batteries since at least 1985, when it was mixed with NaI salt and employed as the separator for an all-solid-state Na– MoS3 battery at 70–98 °C [47]. A drawback of PEO is its relatively low conductivity (1 × 10−5–1 × 10−6 S cm−1 at room temperature) [45]. PEO has been studied with a variety of sodium salts, such as NaClO4, NaI, NaCF3COO, NaPF6, NaCF3SO3, and NaC2F6NOS2 in an effort to boost its conductivity [47–51]. It has received widespread attention due to its relative stability with solid Na, though it is not recognized as stable against molten sodium because 1) its melting temperature is well below that of sodium, and 2) both the PEO and its additives are reactive with molten sodium. Other candidate polymers studied for use in sodium batteries have primarily been studied as gel-polymer electrolytes. These include, for example,

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PVDF–HFP–PEGDME copolymer [52], PMMA [53, 54], and PAN [55]. However, there is no active research into the use of gel-polymer electrolytes in molten sodium batteries due to the high reactivity of molten sodium with the majority of organic and aqueous liquid electrolytes used in these types of separators. As polymer-based separators continue to be developed for emerging solid-state batteries, opportunities may arise for translation of these materials to molten sodium battery systems. Metal Borohydrides and Higher Borates Metal borohydrides and their related compounds have attracted attention due to the discovery of some compounds exhibiting good ionic conductivities in high temperature polymorphs. Metal borohydrides have further sparked research interest due their apparent stability with solid Li and Na metal, though their stability with molten sodium has not yet been determined. To date metal borohydrides are relatively unexplored, and current work has largely focused on Li+-ion conductors as solid electrolytes in all-solid-state batteries, but there has been some work on Na+-ion conductors as well. Metal borohydrides achieve their good conductivities by what has been dubbed the “paddle wheel” mechanism, in which the borohydride anion, (BH4)−, is orientationally disordered within the crystal lattice. This allows the anion complex to rotate to accommodate fast alkali cation (M+) ion conduction [56]. Closo-polyborate salts were identified as possible solid electrolytes with the discovery of superionic Na+ conductivity in sodium dodecahydro-closo-dodecaborate (Na2B12H12) by Udovic et al. in 2014 [57]. At room temperature Na2B12H12 forms a monoclinic phase with B12H122− polyanions with ordered orientation. They reported the material to undergo a phase transition from ordered monoclinic to disordered bcc phase, shown in Fig. 6, enabling the rotation of the polyanion complexes in a manner similar to the “paddle wheel” mechanism of ion conduction [58]. As such

Fig. 6  Crystal structure of disordered cubic Na2B12H12 (cations omitted). Adapted with permission from Ref. [58] Copyright 2017 American Chemical Society

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the high temperature phase of NaB12H12 exhibits very high conductivities of 1 × 10−1 S cm−1 at 540–573 K (267–300 °C). Na2B12H12 displayed hysteretic behavior in its conductivity on heating and cooling, leading researchers to focus on increasing the conductivity by exploration of other closo-polyborate salts and stabilizing the high conductivity disordered phase at lower temperatures [57]. High conductivities of 1 × 10−1 S cm−1 at 400 K (127 °C) and 3 × 10−2 S cm−1 at 297 K (24 °C) have been reported in, respectively, NaCB11H12 and NaCB9H10 [59]. Similar to glass-ceramic materials, mechanical milling was found to stabilize the high temperature, high conductivity phase at lower temperatures [60]. Only recently has work moved beyond fundamental improvement of the conductivity of metal borohydride and closo-polyborate salt electrolytes into real battery applications. Several potential  drawbacks of these salt electrolytes have been identified, however, limiting their application to all-solid-state batteries. NaBH4 has been identified as reacting violently with molten sulfur at temperatures as low as 190 °C to form pure H2 gas [56]. Additionally, metal borohydrides and higher borates are extremely hygroscopic. Some of the precursor materials are also air-sensitive, placing stringent restrictions on manufacturing of these materials. Lastly, it was found that Na2(B12H12)0.5(B10H10)0.5 electrolyte decomposed above 3.2 V vs Na/Na+, which may indicate that these types of electrolytes cannot be used with a variety of battery cathodes for high voltage battery systems [61, 62].

2.3  Cathodes Cathode chemistries vary significantly among molten sodium batteries, and tailoring these materials provides excellent opportunities to innovate new sodium battery technologies. In addition to providing desired battery voltages and usable energy storage capacities, any new cathode chemistry must be compatible with the solid-­ state separators, should be cost-effective, and should minimize or eliminate any hazardous reactions, especially in the event of a separator failure (exposure to molten sodium). One of the most desirable traits of these new chemistries is to enable operation at temperatures considerably lower than the 270–350 °C common to traditional molten sodium batteries. Sulfur The Na–S battery was first patented by Ford Motor Company in 1968 and represents one of the first battery systems to use the molten Na anode. The Na–S battery is simple in its construction, consisting of a molten Na anode, a BASE separator, and a molten sulfur cathode. During discharge, the Na is oxidized to Na+ which migrates through the ionically conductive separator to react with the sulfur being reduced at the cathode to form molten polysulfide Na2S5, which can be further reduced to lower order polysulfides as shown in the reaction:

72 xS + 2Na ↔ Na2Sx (3 ≤ x ≤ 5)

E. D. Spoerke et al. Ecell ~ 2.08 V at 350 °C

In theory the battery can be discharged to Na2S, but in practice both Na2S2 and Na2S are solid at the battery operating temperatures of 350 °C, and the formation of these solid products dramatically increases the resistance to the battery so as to prohibit further discharge. This effectively limits the capacity of the battery to about half its theoretical value. Since its initial discovery, the majority of research on this system has been toward its practical development. The molten sodium, sulfur and polysulfides, and their vapors are highly corrosive, and great effort has been made toward the development of housing materials, their coatings, and sealing materials that are resistant to such corrosion. In the case of catastrophic battery failure, a violent reaction between the molten sodium and the sulfur/polysulfide cathode takes place, further requiring adequate materials and engineering controls. This risk of catastrophic failure has restricted Na–S batteries to utility applications despite the original intent in the 1960s to develop the Na–S battery for vehicles. A further restriction on the Na–S battery applications is the ionic and electronic resistance of the sulfur catholyte. The material’s poor electrical conductivity restricts the battery to low-power applications such as load-leveling and emergency power distribution [63, 64]. Despite these limitations, the Na–S battery has seen over 560 MW (4000 MWh) of storage deployed worldwide, and some of these installations have demonstrated over a decade of reliable cycling [19, 65]. However, recent fires demonstrate the risk of this technology, despite many decades of development [64]. Current research of the Na–S battery has focused on intermediate and room temperature operation to improve battery capacity and safety [66, 67]. Sodium Metal Halides ZEBRA/Na–NiCl2 Closely related to the Na–S system is the sodium metal halide battery, which uses solid or semisolid metal halides as the active cathode material. The most common form of the sodium metal halide battery is the Na–NiCl2 battery, often referred to as a ZEBRA battery (for Zeolite Battery Research Africa, or Zero-Emission Battery Research Activities), as an acknowledgement of the South African origins of its development. This system was originally conceived as a high temperature battery that shared the same type of BASE ion conductor used by the Na–S batteries. In Na–NiCl2 batteries, sodium ions are transported through the oxide membrane from the anode to the cathode during discharge, reducing NiCl2 to Ni via migration of sodium ions in a NaAlCl4 molten salt electrolyte, as shown in the reaction: NiCl2 + 2Na ↔ 2NaCl + Ni(s)

Ecell ~ 2.58 V at 300 °C

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The use of solid or semisolid cathodes makes Na–NiCl2 batteries intrinsically safer and less corrosive than Na–S batteries. Furthermore, the system chemistry allows for assembly of the battery in the discharge state which allows manufacturers to avoid handling sodium metal. Often Na-NiCl2 batteries contain additions of FeCl2 which undergoes the same electrochemical reactions as NiCl2 at the lower voltage of 2.35  V.  This increases the power response of the battery. Pure FeCl2 is rarely used, as FeCl3 can form during overcharge, and reaction between the soluble Fe3+ and the BASE can cause degradation of the solid electrolyte [12]. The molten salt electrolyte, NaAlCl4, also provides some degree of protection from overdischarge in the battery by the formation of Al metal, and serves to protect the battery in the case of catastrophic failure. Contact between the electrolyte and the molten salt result, not in a violent reaction like in the case of Na–S batteries, but instead in the formation of harmless products NaCl and Al. The high voltage of Na–NiCl2 batteries compared to Na–S batteries further helps boost the system energy density. Nevertheless, the Na–NiCl2 battery has yet to achieve widespread use as there remain a number of challenges in its development. Further improvement in power, reliability, and cost are necessary. Particular focus has been placed on reducing the amount of Ni in the battery, which is expensive and weighty, and on reducing the operating temperature to extend the system lifetime. Excess Ni is used in the battery cathode to ensure a good electrical pathway as well as a high surface area for faster reaction kinetics to cycle NiCl2, and thus improve the power performance of the cathode [68]. Serendipitously, it has been found that in reduced Ni batteries, a lower operating temperature is actually required to improve the cyclability of the battery [68]. Na–NiCl2 batteries have thus been demonstrated with Ni contents reduced by as much as 40% with stable cycling up to 150 cycles at 190 °C [8]. Research into reducing the operating temperature of the Na–NiCl2 battery has also allowed the replacement of BASE with higher conductivity NaSICON [69, 70]. Research and development of the Na–NiCl2 will continue to optimize individual parameters to make the system more robust and less expensive to operate and maintain. Hybrid S/NiCl2 Other areas of research on Na–NiCl2 batteries include the use of hybrid, or mixed catholytes to boost the energy density of the system. One example is a hybrid Na–S/ NiCl2 battery which added Na2S to the traditional cathode materials and cycled under the normal operating conditions of a Na–NiCl2 battery. Unlike in a Na–S battery in which the formation of solid low-order polysulfides (Na2Sx, 1 ≤ x ≤ 3) prevents further discharge of the battery, in the presence of molten NaAlCl4 the lower order polysulfides are able to be cycled, boosting the energy density of the system. Despite an initial drop in capacity, likely due to the reaction between Na2S and Ni to form nickel sulfides, the battery exhibited good cyclability in retaining 95% of its capacity. The hybrid Na–S/NiCl2 battery is still in its infancy, having demonstrated only 60 cycles, but the good performance shows promise for this line of work [71].

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Other Metal Halides A wide variety of other metal halides have been screened for use as battery cathodes beyond NiCl2 and FeCl2. Transition metal chlorides such as copper, manganese, chromium, aluminum, silver, and titanium proved to be soluble in the electrolyte and therefore not promising for use in a metal halide battery. Molybdenum, cobalt and zinc chlorides, as well as tin(IV) iodide demonstrated insolubility in the molten electrolyte and are therefore considered to be more promising as cathode materials, though further research is needed to demonstrate the electrochemical effects of cycling the materials in a battery [72–75]. To further enhance the capabilities of the screening process, Zhu et al. developed mathematical formulae to simulate and analyze various molten sodium battery chemistries. In their work they simulated the performance of copper iodide with an iodide salt as the catholyte and high ionic conductivity NaSICON as the solid electrolyte in a low-temperature molten sodium battery [76]. The application of these computer models will help to rapidly identify such promising materials. Fully Molten Salts Inorganic molten salts have recently been explored as a redox-active catholyte for molten sodium batteries. In these systems, the redox-active species is fully dissolved in the molten salt electrolyte to form a redox-active catholyte: an inorganic analogue to the redox-active catholytes in flow batteries. The molten salt catholyte is separated from the sodium anode by a ceramic separator. The use of a catholyte, similar to the molten polysulfides in Na–S batteries, has the advantage of improving redox kinetics over a solid cathode to lower the overall resistance of the battery. An intermediate temperature (120–180 °C) Na–NaI battery has been demonstrated with a NaSICON ceramic in which iodide is the redox-active species per the reaction: [22]. 2Na + I3− ↔ 2Na+ + 3I−

Ecell ~ 3.24 V at 120–180 °C

This system was demonstrated in a scalable design up to 10 Ah. It was further demonstrated to have excellent safety as, like the Na–NiCl2 battery, catastrophic failure leading to contact between the Na and the NaI–AlCl3 catholyte leads to the formation of NaCl and Al [22, 77]. Research in this area is new, and future work is looking to further lower the temperature of the system, optimize low-temperature solid/liquid interfaces, and improve the conductivity of the ceramic separator.

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Ionic Liquids Ionic liquids are starting to attract attention for a wide variety of battery applications, including molten sodium batteries. There are, it appears, two approaches to using ionic liquids. The first is use of the ionic liquid as a traditional electrolyte, as was demonstrated by Peled et al. in a Na-air battery with a molten sodium anode, glass-fiber, and Celgard separators soaked in ionic liquid electrolyte (0.5 M sodium triflate (NaTf)  +  0.1  M Na2SO4 in 1-butyl-1-methylpyrrolidinium bis(trifluoromethanesulfonyl)imide (PYR14TFSI)), and dry air as the cathode. The battery was operated at 105 °C, just above the melting point of sodium. Though an interesting demonstration, the battery was not able to be charged at low current densities owing to a high rate of self-discharge [78]. Another approach is to use the ionic liquid as a catholyte, in which the redox-­ active species are fully dissolved in the ionic liquid, and to protect the molten sodium anode with a solid electrolyte separator. This was demonstrated by Xue et al. [79] in which FeCl3 and NaAlCl4 were dissolved in ionic liquid ethylmethylimidazolium chloride (EMICl) to form EMIFeCl4–NaAlCl4 in which the redox-­ active species were Fe2+/Fe3+ per the reaction: Ecell ~ 3.25 V at 180 °C

Na +  EMI +   FeCl −4  ↔  EMI +   FeCl3−  + NaCl EMICl suffers from alkali cation trapping and poor ionic conductivity at temperatures below 100  °C.  This issue was overcome by operation of the battery at 180 °C, and the use of NaSICON as a separator for good ionic conduction at the intermediate temperature. One primary disadvantage of the demonstrated system was the precipitation of solid NaCl on discharge. This problem, however, can be overcome by the choice of different redox-active species such as, for example, the Fe2Cl7− anion. Aqueous Cathodes An interesting area of research is aqueous cathodes and catholytes. As the melting point of sodium is only a few degrees below the boiling point of water, batteries with aqueous cathodes rely on modification of the anode to lower the melting point of sodium substantially below water’s boiling point. Work at the Pacific Northwest National Laboratory demonstrated that the sodium metal anode can be successfully alloyed with cesium which, depending on the alloy composition, can create a fully liquid sodium anode at room temperature [10]. This novel anode alloy was utilized

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in an aqueous hybrid flow battery with a water-resistant BASE separator and aqueous vanadium flow catholyte [80]. The vanadium catholyte was cycled per the reaction VO2+ + 2H+ + e− ↔ VO2+ + H2O

Ecell ~ 3.7 V at 25 °C

cycling in the V4+/5+ regime to maintain a high cyclability of the cell. The choice of BASE, which has low room temperature conductivity and is known to be unstable in water, required ultra-low current densities of 0.006 mA cm−2 to operate the cells, and ultimately led to poor performance and battery degradation. The low-­ temperature operation of aqueous systems allows substantial flexibility in separator choice, from high conductivity NaSICON ceramic to low-cost polymer separators. The variety of aqueous catholytes available, such as sodium ferrocyanide, aqueous polysulfide, and aqueous air catholytes, shows great promise for these low-­ temperature molten sodium battery systems as well [81–83]. The biggest challenge to the use of these aqueous systems relates to the strong reactivity of sodium and its lower-melting alloys with water. Failure of seals or the separator itself could lead to potentially dangerous side reactions in such a system. Air (O2) Air cathodes use oxygen (O2) in the air as a cathode material. The oxidation and reduction of oxygen is typically mediated by catalysts, often supported on a carbon electrode. This arrangement offers a clear advantage in that the active material does not have to be stored in the battery, dramatically increasing the energy density of the system. The voltage of the air cathode is highly dependent on which reactions occur and the final discharge product. Reactions that may occur are: Na + O2 + e− ↔ NaO2 Ecell ~ 2.26 V 2Na + O2 + 2e− ↔ Na2O2 Ecell ~ 2.33 V 4Na + O2 + 4e− ↔ 2Na2O Ecell ~ 1.95 V

Which reactions are favored are highly dependent on the choice of electrolyte and catalyst. The use of a high temperature air cathode has advantages over its room temperature analogue in that it can be operated in atmosphere, rather than a pure O2 environment, with minimal ingression of water that may react with the sodium anode or cause undesirable side reactions during cycling of the O2. Furthermore, the use of a molten sodium anode will limit dendrite formation in the battery. Despite these advantages, to date there has been limited work coupling air cathodes with a molten sodium cathode. Peled et  al. pioneered a molten sodium-oxygen battery, operated at 105 °C, with a PEO–NaTf electrolyte [78, 84]. The low conductivity of the polymeric electrolyte combined with the sluggish redox kinetics of the O2

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couple restricted the battery to low current densities, but achieved reasonable cyclability for a first demonstration. Future work optimizing the conductivity of the solid electrolyte separator could enhance the power density and lifetime of the system. Molten Metals Due to the restrictions that the solid electrolyte places on cost, manufacturability, and achievable power density (due to its low conductivity), there has been substantial effort to try to eliminate the use of a separator in molten sodium batteries. Work has recently been undertaken in the area of liquid metal batteries, in which a molten sodium anode is separated from a molten metal cathode by differences in density and immiscibility between the anode, molten salt electrolyte, and cathode, without the need for a solid electrolyte. A comprehensive history of the liquid metal battery can be found elsewhere [85]. Many different molten anodes and cathodes have been tested within this system, but with molten sodium anodes, cathodes to date have been restricted to Bi, eutectic Pb–Bi alloys, and Zn. These cathodes have been tested using eutectic NaF–NaCl–NaI, NaI–NaOH or NaI–NaOH–NaNH2, and NaCl– CaCl2 molten salt electrolytes [85–88]. Na–Hg and Na–Sb systems have also been proposed [85]. A common disadvantage of sodium metal liquid batteries is the high solubility of sodium in the molten salts at typical operating temperatures, though this can be partially mitigated by the choice of molten salt electrolyte [89]. This high solubility results in high self-discharge of the battery, which must be counteracted by high current density operation. Overall the high self-discharge of the

Fig. 7 (a) Tubular Na–S battery cell, adapted from image by NASA John Glenn Research Center/ Public domain. (b) Planar Na–NiCl2 battery cell, expanded view. Reprinted with permission from Ref. [90]

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battery has the effect of raising the operating costs of the battery. The combination of the high self-discharge rate and low voltage in the case of the Na–Bi couple (0.74 V) has led research in liquid metal batteries to other molten anode materials beyond sodium.

3  Battery Design Although the focus of this chapter is on the materials chemistry of molten sodium batteries, it is worth briefly discussing the design of these systems, as design can have substantial impact on the materials chemistry and battery performance overall. There are multiple configurations for these battery systems, each appropriate for different battery chemistries, scales, and applications of interest. Figure  7 shows examples of both a concentric tubular configuration and a planar stack configuration for these batteries. Although the basic electrochemistry of the system is not generally expected to be significantly affected by the configuration, it is important to acknowledge that the form of the battery and the material behaviors are connected. One of the principle differences between these two primary designs relates to how the molten constituents remain in contact with the solid electrolyte during charge and discharge cycles. The tubular system relies on gravity and wicking of the molten materials along the sides of the tubular separator, and changes in volume of sodium or catholyte are accommodated by the free volume inside and outside the separator. In contrast, the planar design often requires the planar components be capable of “flexing” under compression to allow for intimate contact of molten species with the separator while accommodating changes in volume during charge and discharge. Depending on the specific materials chemistry, these design parameters can significantly influence battery performance. It was demonstrated that for Na– NiCl2 batteries, changing from a tubular to a planar design allowed for lower operating temperatures, improved energy density, and extended battery cycle life [90]. Form has an impact on materials chemistry through the battery cost as well. Expensive configurations, in turn, require less expensive cathodes and separators to balance the cost of the system. Additional components of the battery assembly must also be carefully selected, ranging from the battery housing to the battery seals. High temperature operation places restraints on the types of materials that can be used to house and seal the battery, as does the use of highly corrosive materials or volatile species [19, 91]. Finally, it must be mentioned that there are numerous practical considerations for how the system will be assembled and sealed when developing molten sodium battery chemistries. For example, it is possible to assemble cells in the fully discharged state, meaning that sodium metal need not be handled in large quantities during assembly; sodium is generated during an initial charge state. While this is a desirable processing approach, it does introduce complications of providing free volume for the sodium generated on charge and maintaining suitable electrical contact

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between the current collector and separator during initial charge. Each of these factors will vary as a function of the intended scale, application, and cost target of the system being designed.

4  Current Battery Deployments Despite the reliance on higher temperature operation, as of early 2020, both Na–S and ZEBRA batteries have been deployed in communities around the world for a range of grid-scale applications, and it is anticipated that additional battery developers in Germany (Fraunhofer Institute for Ceramic Technologies and Systems (IKTS)) and China (Chilwee Group) may be aiming to expand commercial molten salt (ZEBRA-based) battery utilization in coming years. Na–S batteries have, to date, largely been deployed by NGK Insulators. As mentioned earlier, NGK’s “NAS” batteries have been deployed in over 200 locations to enable over 560  MW/4000  MWh of storage. These batteries have a functional energy density around 300-400 Wh/L and are expected to run 300 cycles per year for 15 years. Although they operate near 300 °C, they can be used in a range of ambient temperatures ranging from −20 °C to +40 °C. They have found uses across industrial, commercial, and residential application space, enabling renewable integration, investment deferral and ancillary services, and powering remote communities or microgrids. The systems are scalable in size, evidenced by the installation of smaller systems, such as a 1 MW microgrid support system on Catalina Island, CA (USA), contrasted with a 50 MW/300 MWh system supporting a solar installation in Fukuoka, Kyusyu, (Japan) [6]. Na–NiCl2 systems have been primarily deployed by FZSoNick, employing grid-­ scale batteries suitable for frequency regulation, load shifting, peak shaving, backup power, and renewables integration [3]. These systems which operate near 270 °C have a functional energy density around 150–190 Wh/L with an expected lifetime of around 20 years and are also suitable for operation in variable ambient temperatures (−40  °C to +60  °C). They categorize their technology space into energy backup, energy storage, and additionally in mobile applications (vehicles). Much of the stationary storage applications have focused on backup power for telecommunications, public transportation, and remote site applications. At present, FZSoNick has deployed roughly 100 MWh of energy storage with approximately 2 MAh of backup storage capacity, most heavily focused around telecommunications backup. An additional 14 MWh of storage is used for to enable renewables integration, microgrid applications, and grid services (e.g., grid balancing, voltage regulation). In addition, these batteries are enabling vehicle electrification of buses, light commercial vehicles, and industrial machinery. That these batteries are able to provide the combination of performance and safety needed for these mobile applications is an important consideration for their continued development and implementation in the growing energy storage marketplace.

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5  Conclusion Sodium is a globally abundant, inexpensive, and high energy density material that has an established precedent and promising future as the basis for a family of grid-­ scale batteries. With a moderate melting temperature, molten sodium offers distinct advantages in terms of electrochemical kinetics and flexibility of battery form design. Traditionally elevated operating temperatures of Na–S and Na–NiCl2 batteries, however, contribute to market-limiting costs and may need to be overcome to enable the widespread utility of this technology. Widely varied potential cathode chemistries offer opportunities to lower operating temperatures, increase voltages, extend battery lifetimes, and reduce system costs. In addition, advancements made toward lower temperature, high conductivity separator materials will further drive down operating temperatures and improve both the reliability and performance of sodium batteries. Innovations in the materials chemistries of these batteries will also inform novel cell designs that optimize chemical and electrochemical properties of emerging systems to drive the technology forward. Meanwhile, current deployments of Na–S and Na–NiCl2 battery systems, based on traditional high temperature systems, enable a wide range of grid-scale services around the globe with promise toward continued expansion as technical and market advances drive down costs and create opportunities for these technologies. Active research and forward-thinking development present great promise for a new generation of safe, reliable, long-lived grid-level energy storage technologies. Acknowledgements  The authors were supported through the Energy Storage Program, managed by Dr. Imre Gyuk in the U.S.  Department of Energy Office of Electricity. Sandia National Laboratories is a multimission laboratory managed and operated by National Technology & Engineering Solutions of Sandia, LLC, a wholly owned subsidiary of Honeywell International Inc., for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-NA0003525. This paper describes objective technical results and analysis. Any subjective views or opinions that might be expressed in the paper do not necessarily represent the views of the U.S. Department of Energy or the United States Government.

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76. Zhu H, Bhavaraju S, Kee RJ (2013) Computational model of a sodium-copper-iodide rechargeable battery. Electrochim Acta 112:629–639 77. Percival SJ, Small LJ, Spoerke ED (2018) Electrochemistry of the NaI-AlCl3 molten salt system for use as catholyte in sodium metal batteries. J Electrochem Soc 165(14):A3531–A3536 78. Peled E, Golodnitsky D, Hadar R, Mazor H, Goor M, Burstein L (2013) Challenges and obstacles in the development of sodium–air batteries. J Power Sources 244:771–776 79. Xue L, Tucker TG, Angell CA (2015) Ionic liquid redox catholyte for high energy efficiency, low-cost energy storage. Adv Energy Mater 5(12):1500271 80. Liu C, Shamie JS, Shaw LL, Sprenkle VL (2016) An ambient temperature molten sodium-­ vanadium battery with aqueous flowing catholyte. ACS Appl Mater Interfaces 8:1545–1552 81. Senthilkumar ST, Han J, Park J, Hwang SM, Jeon D, Kim Y (2018) Energy efficient Na-aqueous catholyte redox flow battery. Energy Storage Mater 12:324–330 82. Gross MM, Manthiram A (2019) Development of low-cost sodium-aqueous polysulfide hybrid batteries. Energy Storage Mater 19:346–351 83. Hwang SM, Go W, Yu H, Kim Y (2017) Hybrid Na–air flow batteries using an acidic catholyte: effect of the catholyte pH on the cell performance. J Mater Chem A 5(23):11592–11600 84. Peled E, Golodnitsky D, Mazor H, Goor M, Avshalomov S (2011) Parameter analysis of a practical lithium- and sodium-air electric vehicle battery. J Power Sources 196(16):6835–6840 85. Kim H, Boysen DA, Newhouse JM, Spatocco BL, Chung B, Burke PJ, Bradwell DJ, Jiang K, Tomaszowska AA, Wang K, Wei W, Ortiz LA, Barriga SA, Poizeau SM, Sadoway DR (2013) Liquid metal batteries: past, present, and future. Chem Rev 113(3):2075–2099 86. Spatocco BL, Ouchi T, Lambotte G, Burke PJ, Sadoway DR (2015) Low-temperature molten salt electrolytes for membrane-free sodium metal batteries. J Electrochem Soc 162(14):A2729–A2736 87. Ashour RF, Yin H, Ouchi T, Kelley DH, Sadoway DR (2017) Communication—molten amide-hydroxide-iodide electrolyte for a low-temperature sodium-based liquid metal battery. J Electrochem Soc 164(2):A535–A537 88. Xu J, Kjos OS, Osen KS, Martinez AM, Kongstein OE, Haarberg GM (2016) Na-Zn liquid metal battery. J Power Sources 332:274–280 89. Spatocco BL (2015) Investigation of molten salt electrolytes for low-temperature liquid metal batteries. Doctoral dissertation, Massachusetts Institute of Technology 90. Li G, Lu X, Kim JY, Meinhardt KD, Chang HJ, Canfield NL, Sprenkle VL (2016) Advanced intermediate temperature sodium–nickel chloride batteries with ultra-high energy density. Nat Commun 7:10683 91. Chang H-J, Lu X, Bonnett JF, Canfield NL, Son S, Park Y-C, Jung K, Sprenkle VL, Li G (2017) Development of intermediate temperature sodium nickel chloride rechargeable batteries using conventional polymer sealing technologies. J Power Sources 348:150–157

Polymer Nanocomposites for Ion Transport Christina A. Bauer

Abstract  Directional movement of ions between electrodes in a safe, rapid, and efficient manner requires rational design of functional solid-state electrolytes. Although liquid and polymer electrolytes have resulted in major technological breakthroughs, the next generation structures aim to make green energy solutions more widespread and feasible. Nanostructures offer the ability to construct complex, tailored assemblies that can address the current limitations to ion mobility in solid materials by providing liquid-like movement in the solid state, expanding operational voltage windows, and improving durability and temperature stability. Two types of polymeric nanocomposites will be addressed here: (1) Hybrid materials that consist of organic polymers with inorganic nanoparticles and (2) Porous coordination polymers, including metal organic frameworks and zeolitic imidazolate frameworks. The ability to tailor such porous coordination polymers in all aspects—the metal, the organic linker, and the identity of the guests—make these viable candidates for future green energy solutions. The current progress in applications toward fuel cells (acidic and basic), batteries, supercapacitors, and solar cells will be summarized and described here, along with a perspective on challenges and future directions. Keywords  Nanocomposites · Metal organic frameworks · Ion conductivity · Solid-state electrolytes · Coordination polymers · Batteries · Lithium ion batteries · Supercapacitors · Green energy · Fuel cells · Solar cells · Dye-sensitized solar cells · Anion-exchange fuel cells · Cation-exchange fuel cells

1  Introduction Practical, clean energy storage solutions are of tremendous interest. The threat of an unstable oil market, the need to supply electrical power in remote locations and under extreme conditions, and the desire to increase the efficiency of energy conversion while reducing greenhouse gas output have culminated in a boom in research C. A. Bauer (*) Department of Chemistry, Whittier College, Whittier, CA, USA e-mail: [email protected] © National Technology & Engineering Solutions of Sandia, LLC 2021 M. Alston, T. N. Lambert (eds.), Energy-Sustainable Advanced Materials, https://doi.org/10.1007/978-3-030-57492-5_4

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for the alternative energy fields. Synergistically, this comes at a time when the synthetic ability to make and control various modular hybrid organic/inorganic materials for advanced functionalities is drastically improving. The types of applications that have been studied thus far include supercapacitors, batteries, fuel cells, and solar cells, to name the major ones to be discussed here, alongside other devices such as electrical gas sensors. The requirements have become such that more convenience and mechanical flexibility in the devices, along with greater energy capacity per unit mass or unit volume, with the ability to rapidly charge and discharge for high power operation is needed. Moreover, the goal is to gain more specificity with little need for maintenance, and perhaps most importantly, long-lasting and safe technology. Without the advent of advanced energy solutions, lightweight, modern electronics would not be possible. For example, compare a heavy lead-acid battery to the lithium ion batteries used in laptops today. The advancements in and miniaturization of technology and electronics need to be accompanied by similar development of their power sources. Although major improvements have been made, and are continuously advancing, most current technology is based on liquid organic electrolytes, which have considerable drawbacks, such as potential for leakage, flammability, and narrow temperature ranges of operation. Therefore, solid or elastomeric organic materials are sought after, having the advantage of being generally lightweight, processable, and flexible, with relative ease of manufacturing in many cases. In addition, there is an improvement in robustness and safety when compared to liquids. At the heart of all electrochemical devices is the electrolyte, as depicted in simplified schematics in Fig. 1. Despite differences in construction, they fundamentally utilize connected electrodes separated by an electrically insulating membrane to force the electrons to move through the wire and therefore generate electrical work. Movement of electrons needs to be balanced by the movement of ions between the electrodes. Electrolytes conduct charge by an accompanying movement of ions, and in these devices, they are chosen for their ionic transport capabilities. Electrical work is done directionally, and ion movement in an organized, directional fashion may also be of use to minimize randomness and self-diffusion of ions. Because of the inherent limitations of liquid-phase materials (including both aqueous and organic solvents), these fundamentally do not meet the demands of future devices and require considerable improvement. Ionic liquids can provide an enhancement in terms of conductivity and thermal stability with a much lower vapor pressure at standard conditions, but still do not solve all issues. So, many have shifted their attention toward solid solutions. A common theme presents in the development of these technologies, and each goes through a similar maturation process. The first devices are made with liquid electrolytes, which are simpler and effective (such as sulfuric acid in lead-acid batteries). Once optimization reaches a certain point, ionic liquids are utilized to allow for more robustness and wider operational capability. Finally, the solid state becomes important, even crucial, for first generation commercial devices.

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Fig. 1  Schematic of electrochemical devices involving an external circuit to harness electrical work. The anode is on the left, the cathode on the right in all cases as drawn here, allowing electrons to flow to the right. In Li-ion batteries, the ion intercalates into the cathode upon discharge and intercalates into the anode upon charging. In proton conducting fuel cells, the fuel is supplied at the anode, with the oxygen at the cathode. Protons move to the right. In anion-exchange membrane fuel cells, the fuel is supplied at the anode, with oxygen and water at the cathode. Hydroxide ions move left. Supercapacitors are a symmetrical sandwich, utilizing nanosized conductors in contact with a thin layer of electrolyte, separated by an insulator. A static charge is imposed upon charging, aligning the dipoles in the electrolyte. Upon discharge, electrons flow to the cathode very quickly. In solar cells, light enters through transparent glass and is absorbed by the dye (the sensitizer). This excited dye transfers an electron to the wide-gap semiconductor (commonly TiO2 nanoparticles), which then transfers to the anode. Meanwhile, the redox component allows for replenishing the dye to enable for continued use

Solid electrolytes are highly sought after, but common solids are often not particularly good ion conductors with some well-known exceptions such as α-silver iodide, which conducts silver ions very efficiently [1]. Ionic solids typically work via transport of ions through crystalline point defects (i.e. Schottky) or dislocation defects (i.e. Frenkel). However, these solid materials are lacking in flexibility and are difficult to integrate with the electrodes. Instead, attention has turned to high dielectric constant polymers that coordinate and solvate ions, hence behaving as a solid solvent for these ions, analogous to ions in solution. Both dry and solventincorporated versions have been developed, and each is important in different ways, depending on the temperature and conditions of operation. Polymeric materials offer the ability to form solid films, and can contain amorphous groups that allow for liquid-like movement and solvation of ions, provide robustness and longevity,

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and impose less caustic conditions on the electrodes than liquids. Furthermore, the leakage and flammability issues are thwarted in many cases. This chapter is intended to discuss the main types of technologies that utilize such composite solid electrolytes, and to compare and contrast the needs that may be met with them, as well as how they might be tailored for specific applications. In particular, organic polymers and composites thereof, along with metal organic coordination polymers, have shown considerable promise in this area. Among types of ionic conductors, “mixed” conductors are able to transport both electrons and ions. Those solids with preferential ionic conductivity are often referred to as solid electrolytes or superionic conductors (which implies a high conductivity value, see specific definition below), allowing for the rapid movement of ions through the network. When only one ion moves in response to a field, the material is known as a single-ion conductor, a feature that is often achieved via incorporation of anchored counter-ions in the solid matrix. A detailed description can be found in “Solid Electrolytes: Materials, Properties and Applications.” [2]. To behave solely as an electrolyte in applications, the material must be electrically insulating in order to prevent short-circuiting (i.e. be a dielectric). However, electrodes must possess both ion and electron conducting properties. The focus here is on the electrolyte, but the interested reader is directed to consider other work on polymeric nanocomposites for electrode materials, a rich field of study [3]. Ionic conductivity can be formulated similarly to electronic conductivity, by considering the size, the concentration, and mobility of the charge. In Eq. (1) below, the product of these terms describes the ionic conductivity (σi), where mobility (μ) is defined as the speed of charge movement in the presence of an electric field, Z is the magnitude of the charge of the ion, and n is the concentration of the carriers.

σ i = n·Z ·e·µ

(1)

Ionic conductivity is typically not as high as electronic conductivity, so it can be a limiting factor in device performance. Traditionally in solid electrolytes, hopping of either cations or anions between lattice sites occurs when a potential is introduced. The process is activated, with an associated Arrhenius activation energy (Ea). The defects required for mobility can be intrinsic (Frenkel or Schottky) or extrinsic (introduced by doping). Carrier concentration will vary with temperature for thermally induced intrinsic defects. For example, the α phase of AgI occurs above 147 °C, whereby Ag+ ions exhibit liquid-like movement due to temperature-induced cationic disorder. Ideally, one looks for ion specificity, low activation energies to initiate transport, and high ion conductivities. The possible mechanisms are dependent upon the system and are a continual area of study as new, diverse materials for good electrolyte behavior are developed that may require adjustment to the classic models. Low associated Ea often indicates that a hopping mechanism, such as in the well-known Grotthuss mechanism for hydrated protons (See Fig. 2), is likely vs. a vehicle-type mechanism, whereby charges move along with assistance from another species. These concepts are also useful for describing polymeric electrolytes. In terms of porous coordination polymer frameworks, intrinsic usually refers to

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89

Fig. 2  The vehicle mechanism is described by the movement of an ion (blue) that is transported with another carrier (the vehicle) and requires higher activation energy (0.50–0.90 eV) than the Grotthuss mechanism (0.10–0.40 eV), whereby the ion (specifically a proton) hops from one stationary site to another. The counter diffusion of unprotonated vehicles (e.g. H2O) results in the net transport of protons, so this type of conductivity is a function of the medium (Reprinted with permission from ref. [4])

conductivity by the framework itself, whereas extrinsic conductivity is due to incorporated guests or post-synthetic modifications of the framework [5]. Ionic conductivity (σi) in solids is measured in terms of Ohm−1 cm−1 or S cm−1 (Note that S stands for Siemens, which is the inverse of resistance, having the unit of Ohm−1). When σi is on the order of 10−4 to 10−1 S cm−1 at 300 K, a material is classified as a superionic conductor, and when both σi > 0.1 S cm−1 at 300 K and the activation energy for ion transport is small (about 0.1 eV), materials are classified as advanced superionic conductors. The most famous example of an advanced superionic conductor-solid electrolyte is RbAg4I5 where σi > 0.25 S cm−1 and the electrical conductivity, σe, is very small ~10−9 S cm−1 at 300 K. Another term for a solid electrolyte is a “fast ion conductor,” which is considered to be useful when the conductivity is at or above 1 × 10−4 S cm−1 for lithium ions [6], and 1 × 10−2 S cm−1 for protons [7]. The bulk ionic conductivity of a material reflects transport of both anions and cations as the sum of all mobile ion conductivities. Often, for practical devices the effective conductivity of a specific ion in the material is of paramount interest. This is defined as the transference number for the specific ion, formulated as the fraction of the ionic current carried by this ion divided by the total ion current. The closer this transference number is to unity, the more specificity for movement of that particular ion in the material (approaching 1 for single-ion conducting materials). As protons are usually able to move quicker than other species, transference numbers are not typically reported for proton conduction and are assumed to be very close to 1. For other ions, however, they can be very important, in particular for electrolytes

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used in lithium ion batteries. To date, H+, OH−, and Li+ represent the vast majority of mobile ions studied in hybrid solid ion conductors, and the majority of this Chapter will focus on these alongside a brief survey of related Na+ and Mg2+conducting materials.

2  Polymer Nanocomposites The use of polymers allows for robustness, safety, and flexibility, yet the freedom of ion movement can be limited with micro-aggregation or crystallization, leading to lower-than-ideal conductivities. It is typically necessary for the polymer phase to be amorphous in order for ions to move (although a report in 2001 showed ion conductivity in a crystalline polymer) [8]. How can the polymers be manipulated to become more efficient ion conductors? Also, how can modern electrochemical devices take advantage of the many attractive properties presented by polymers, but reinforce them mechanically? The answer may lie in an integrated mixture of organic and inorganic entities. The flexibility afforded by organic polymers and the strength provided by the inorganic portions allows for benefits of both. Two broad types of ion conductors that will be discussed here involve (1) mobility through a nanoscale polymeric composite that is capable of solvating ions or has been functionalized with anchored moieties that can do so, and (2) mobility along an ion conductor, organized in particular geometries with polymeric nanomaterials. Advanced strategies are now being applied to maintain an amorphous composition at a range of temperatures, hence a high conductivity under technologically useful conditions.

2.1  Organic Polymers/Inorganic Nanoparticles Organic polymers are popular as they have revolutionized much of the world we live in today, in part due to their relative ease of processing; they can be molded and set with ease, and can be very chemically and thermally stable [9]. They may also be tailored to allow for liquid-like behavior within a solid form as advances in polymer chemistry allow for control of size, organization, and functionality of these materials. Most neutral polymers are poor ionic conductors on their own (σi 10 water/unit cell Asymmetric, homochiral

Highly oriented thin films

Nonporous, vehicle mechanism

Impregnated ionic liquid

[52] (continued)

[51]

[48] [49] [50]

[46] [47]

[45]

[29]

[44]

Functionalization with ethylene [41] glycol Guest-induced SCSC, proton [42] released MOF/Nafion®/phytic acid hybrid [43]

Isostructural series, alter pKa of [39] R 1-D metal organic nanotubes [40]

Conductivity (S cm−1)/Ea/Temp/RH Comment

Type

Metal Organic linker Proton conduction Al or Fe BDC-R

Table 1  Summary of MOFs used for proton conduction

Polymer Nanocomposites for Ion Transport 99

5-sulfoisophthalate [Ru(4,4′-dcbpy)3]4− Benzene-1,2,4,5-tetramethylenePA

In La La

Oxalate Tris(hydrogen phosphonate)

2,4,6-trihydroxy-1,3,5-benzenetrisulfonate 2,5-dioxido-1,4-BDC

3-methyl-2-(pyridin-4-ylmethylamino) butanoic acid] 2,5-dioxidoBDC 1,2,4-triazole, orthophosphates Bis(4-carboxyphenyl)imidazolium

adp, oxalate

Mn/Cr Na

Na Ni

Zn

Zn

Zn Zn Zn

inh.

2,5-dihydroxy-1,4-benzoquinone

Mn

PSM, guest

PSM inh. inh.

inh.

Guest PSM

10−2, 0.63 eV, R.T., 98%RH

1.1 × 10−3, 0.23 eV, R.T., 96% RH 2.1 × 10−2, 0.21 eV, 358 K, 90% RH 5 × 10−4, 0.34 eV, 423 K, Anh. 2.2 × 10−2, 0.12–0.20 eV, 353 K, 95%RH 4.45 × 10−5, 0.34 eV, 304 K, 98% RH 4.3 × 10−9, 419 K, anh. >10−4, 0.60 eV, 423 K, anh. 2.3 × 10−3, 0.22 eV, R.T. 95% RH,

2.7 × 10−4, 0.47 eV, 363 K, 98% RH 4 × 10−5, 0.26 eV, 298 K, 95%RH

inh.

Guest PSM

6 × 10−3, 0.17 eV, 358 K, 95% RH

[60]

[59]

Triazole guests pH dependence, soaked in sulfuric acid Water chains confined in hydrophobic pores Post-grafted histamine Anisotropic conductivity Low conductivity until fifth water introduced Rational design

[70]

[67] [68] [69]

[66]

[64] [65]

Water-induced structure change, [61] reversible Ferromagnetic [62] Isomorphous ligand replacement [63]

Anionic, single crystal Visualization via luminescence Utilizes phosphonic acid group, resists swelling Particle size dependence: Bigger = better conductivity Novel water aggregates

[56] [57] [58]

[55]

4.7 × 10−4 S cm−1, 0.70 eV, >348 K, 95% RH >10−3, 0.31 eV, 298 K, 40% RH 5.5 × 10−7, 313 K, 95% RH 4 × 10−3, 0.32 eV, 333 K, 98% RH Channel-assisted, magnetic

Reference [53] [54]

Conductivity (S cm−1)/Ea/Temp/RH Comment 1.3 × 10−3, 0.37 eV, R.T., 98%RH 1-D, no acid groups ∼10−4 S cm−1, R.T., 65% RH Low humidity

inh.

inh. inh. inh.

inh.

Type inh. inh.

La,Ce, Pr, Nd, 1,2,4,5-tetrakis(phosphonomethyl)benzene Sm, Eu, Gd Mg, Sr 2,2′,6,6′-tetracarboxybiphenyl

Gd, Dy

Organic linker Oxalate NR3CH2COOH, R = methyl, ethyl, or n-butyl; oxalate Mucic acid, oxalic acid

Metal Fe Fe/Cr; Mn/Cr

Table 1 (continued)

100 C. A. Bauer

Guest inh. inh. inh.

Triazole

Benzimidazole, and orthophosphate Oxalate

1H-1,2,3-triazole, BDC

NH4+, phosphate NH4+, phosphate

(CH3)2NH2, oxalate

1H-pyrazole, 1H-1,2,4-triazole

Zn

Zn Zn

Zn

Zr Zr

Zr/Li

Zr

Zn Zn

Guest Guest

inh. Guest

6.0 × 10−5, 0.38 eV, R.T. 99.9% RH, 0.145, 353 K, 100% RH 1.16 × 10−2, 0.20 eV, 313 K, 100% RH 8 × 10−5, 0.19 eV, 300 K, 99%RH 2.3 × 10−8, 0.70 eV, 298 K, 98% RH N/A 9.8 × 10−4, 333 K, 33%RH

1.3 × 10−3, 0.40 eV, 393 K, anh. 1 × 10−4, 0.130 eV, anh. 4.2 × 10−2 98% RH, 423 K 2.65 × 10−4, 0.18 eV, up to 393 K, anh. 1.45 × 10−3, 0.26 eV, 453 K, anh. 1.45 × 10−2, 0.19 eV, 363 K, 95% RH, 1.1 × 10−5 503 K, anh. 3.9 × 10−5, 0.64 eV, 67% RH, 290 K >10−1, 0.10–0.28 eV, 358 K, 90% RH

3.5 × 10−5, 0.17 eV, 298 K 98% RH, 8.4 × 10−3, 0.25 eV, 300 K, 98% RH 4.6 × 10−3, 0.53 eV, 423 K, anh.

[82] [83]

In situ incorporated OH− anions Salt inclusion, first MOF for OH− cond. Used for OH− exchange Ionic liquid, low RH

[84] [85]

[80] [81]

[79]

[78]

[77]

[75] [76]

[10]

[5] [74]

[73]

[72]

[71]

Anionic stripping PSM with KOH

Anionic stripping

Insulator-conductor transition, dense Superprotonic

Anionic framework High temp range

Nonporous, encapsulate H+ carriers in defect sites Templated Works in hydrous and anhydrous conditions SCSC, large temp. range

[(CH3)2NH2]+ guests

Interlayer lined with water

Type of conductivity: Guest (extrinsic), PSM (extrinsic), inh. = native framework via included functionalities (inherent or intrinsic) RH = relative humidity, anh. = anhydrous, bipy = bipyridyl, terephthalic acid = BDC = 1,4-benzene dicarboxylic acid, meim = methylimidazole, BTC = 1,3,5 benzenetricarboxylic acid, PA = phosphonic acid, adp = adipamide, SCSC = single-crystal-to-single-crystal transition

2-meim(ZIF-8) 2-meim(ZIF-8)

Ni Zn

PSM PSM

Fe Ni

1-amino-BDC 1,4-bis(pyrazol-4-yl)benzene-4-X with X = H, OH, NH2 2-pyrimidinecarboxylate 2-meim(ZIF-8)

PSM

Hydroxide conduction Cr BTC (MIL-100)

PSM

inh. inh.

inh.

Guest

3,3′,4,4′-biphenyltetracarboxylate

Zn

inh.

1,3,5-benzenetriPA

Zn

Polymer Nanocomposites for Ion Transport 101

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treated somewhat like a low-density filler, analogous to the polymer/nanoparticle composites that came before. For example, phytic acid was incorporated into MIL-101, and used as a filler for Nafion®. Enhancement of the thermal and mechanical properties of Nafion® was noted, and conductivities that were 2.8 and 11.0 times higher than that of pristine membrane were measured for some different variations [43]. Additionally, MIL-101 was impregnated with binary ionic liquid to function safely above 100 °C, allowing for anhydrous proton conductivity. The ionic liquid 1-(1-Ethyl-3-imidazolium)propane-3-sulfonate (EIMS, pKa ≈ 6.8) is a zwitterionic liquid, which has both a cation and anion that are tethered together and cannot migrate along potential gradients, and thus may favor only proton conduction in MIL 101 [44]. MOF frameworks can be anionic, cationic, or neutral, depending on their components. In the case where a neutral MOF demonstrates appreciable inherent proton conductivity, open metal sites will often behave as Lewis acids to facilitate ion transport. An example of conductivity of a neutral MOF can be found in Cu-based HKUST-1, which contains accessible Cu coordination sites. These open metal sites are acidic enough to react with coordinated water molecules and increase the number of protons available for transport within the methanol guests. This conductivity does require a polar solvent to solvate the proton, and drops significantly in acetonitrile, becoming negligible in hexane. As the solvents can be exchanged within the pores, the conductivity can be readily tuned [48]. In other neutral MOFs, the organic linker may be functionalized to allow for a charge, whereby acidic groups within the framework structure can be used to coordinate to and transport protons. Kitagawa reported on the rational design of highly proton-conductive MOFs. In that work, three specific types of changes were made. In the simplest of the three, acid groups could be placed into the pores of frameworks to introduce them directly as counter-ions such as NH4+, H3O+, and HSO4−. The second way was to put acid groups on frameworks, where the protons are provided from them initially. The third was to incorporate acidic molecules into voids (See Fig.  5, left). Inspired by this classification, the following was used here in

Fig. 5  Left shows an idealized pathway for proton movement through continuous pores of a MOF (Reprinted with permission from ref. [25]). Right shows an example of a nonporous coordination polymer, whereby phosphoric acid is in the defect vacancies (Reprinted with permission from ref. [73])

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Table 1: PSM refers to post-synthetic modification for conductivity, inh. refers to inherent conductivity of the unmodified framework, and guests refers to use of incorporated guests as charge carriers. Kitagawa demonstrated a combination of two of the concepts by introducing NH4+ ions using an anionic framework and putting carboxyl end groups of adipic acid into honeycomb-shaped voids. This design resulted in superprotonic conductivity of 10−2 S cm−1 at ambient temperature, comparable to Nafion®. The activation energy (Ea) for this MOF was found to be 0.63 eV. The mechanism of proton conduction is expected to be influenced by both the Grotthuss and vehicle-type mechanisms, whereby there is direct diffusion of protons. As the conductivity remains high in spite of this activation energy, one might expect there to be a high concentration of charge carriers [70]. Kitagawa et  al. also later reported an isostructural series, whereby changes to functional group ligands on the organic linkers were done in a systematic way to observe effects on proton transport [39]. Here, a series of iron- or aluminum-based MOFs were made with alterations to the ligand’s acidity. It was found that the smaller the pKa, the better the proton transport properties, owing to the increased number of protons. In another study utilizing a Ni-based MOF, the acidified MOF exhibits a proton conductivity of 2.2 × 10−2 S cm−1 at 80 °C and 95% RH at pH 1.8 with low activation energy [65]. Since proton conductivity is determined partly by the amount of protonic charge carriers, it is advantageous to develop methods to increase the proton concentration in a material. A strong acid can be fully ionized in an aqueous solution and H+ can be directly supplied to the MOF of interest, which requires that the framework should be stable under such acidic conditions. However, in contrast, there are no strongly acidic groups in the one-dimensional coordination polymer ferrous oxalate dihydrate, which showed proton conductivity on the order of 1.3 × 10−3 S cm−1 at ambient temperatures, suggesting other mechanisms are at play [53]. The permanent and tunable porosity of MOFs has been exploited in a number of ways. As the structures remain crystalline upon evacuation of the pores, these empty spaces can be filled with guest molecules. In 2009, it was first shown that a MOF could be incorporated with guests to exhibit anhydrous proton conductivity. Na3(2,4,6-trihydroxy-1,3,5-benzenetrisulfonate), which has regular one-­ dimensional pores that are lined with sulfonate groups, was modulated by the controlled loading of 1H-1,2,4-triazole (Tz) guests within the pores, reaching 5 × 10−4 S cm−1 at 150 °C. Furthermore, this showed to be gas tight, an improvement over polymer membranes [64]. Another group demonstrated the enhanced proton conductivity in a 3D MOF by the cooperation of guest [(CH3)2NH2]+ cations with water molecules and the carboxylates in the host frameworks. The low activation energy of 0.25 eV suggests a Grotthuss mechanism, and is attributed to high proton carrier concentration and adsorbed water molecules [72]. PSM of MOFs has been used to introduce histidine, a protic amino acid in biologically inspired membranes. Although the observed anhydrous conductivity is not high (4.3  ×  10−9  S  cm−1 at 146  °C), the ion-conductive range of temperatures is largely shifted from 40–80  °C to 120–146  °C compared with the bulk histamine [67]. A more recent paper shows 1000-fold enhancement in proton conductivity of

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two isomorphous, dense metal-terephthalate (m = Mg, Cd) MOFs, and one related Nd MOF using post-synthetically anchored proton transporters. Zwitterionic 4-­pyridinol was included in the original synthesis, which provides a coordinating proton source and the MOF was then additionally loaded with ethylene glycol molecules via partial exchange of the pyridinol groups, allowing for conductivity of protons via hydrogen-bonding type interactions. Not only did the hydrogen-­bonding environment result in increased conductivities (up to 10−3 S cm−1 for Cd), the Ea was one of the lowest reported thus far, at 0.11 eV [41]. When designing a MOF that will conduct protons “as-is,” it is important to consider what the overall structure is and, specifically, the orientation of the acidic or basic groups, which ideally are pointing toward the continuous porous connections. Furthermore, these can provide water-assisted conductivity via channels for water to travel through. MOFs often have anisotropic packing that can allow for observation of directional transport that is not as readily achieved in standard organic polymer composites. A few Cu-based MOFs serve as examples, including a polyoxymetalate MOF containing a 1-D water channel [50] and an anisotropic, homochiral 2D MOF containing hydrogen-bonded waters [47]. Also, a lanthanide MOF, which has two distinct channels decorated with pendant hydroxyls was studied. One channel has more water per area, and correspondingly, higher transport along this axis was predicted via computation [55]. Also a Zn-based MOF containing imidazolium groups aligned inside the channel along the crystallographic a and b axes showed low conductivity up to the tetrahydrate stage, but once the fifth water molecule is introduced, it shows high proton conductivity with a low activation energy, both of which are comparable to those of Nafion® [69]. Isostructural functionalized metal–organic nanotubes have been synthesized using 5-triazole isophthalic acid with proton conductivity along channels, measured as 5.35 × 10−5 and 3.61 × 10−3 S cm−1 for In and Cd, respectively [40]. Magnetic MOFs were also used to arrange the pores in MOFs. Cu allowed for ferrimagnetic ordering and pores containing water to mediate proton transport via proton exchange along ammonium and hydronium groups [49]. Phosphonates and phosphates are finding application in proton conducting MOFs. Phosphonate groups typically offer high water stability and the potential for ligating as a hydrogen phosphonate, yielding additional acidic pores [59]. The first intrinsic proton conductivity in a MOF was observed in 2012. Here, Zn2+, triazole, and orthophosphates result in a 2D layered structure with a conductivity of >10−4 S cm−1 at 150 °C parallel to the layers with a low activation energy [68]. A water stable carboxylate La MOF containing phosphonic acid groups that resists swelling upon proton incorporation was reported by Shimizu et al. The conductivity was measured to have an activation energy of 0.32 eV, which the authors suggest is still within the Grotthuss range, but indicates less ordering [58]. Also, isomorphous ligand replacement of trisulfonate with bis(hydrogen phosphonate) in a sodium MOF leads to 1.5× enhancement in proton conduction with greater acidity, having pores partially lined with hydrogen phosphonates [63]. Most recently in 2018, a MOF containing zirconium phosphate in one-dimensional anionic chains, which is charge-balanced with NH4+ cations yielded a stable anhydrous proton conductivity

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of 1.45  ×  10−3  S  cm−1 at 180  °C, one of the highest values among reported intermediate-­temperature proton conducting materials. When assembled into a H2/ O2 fuel cell, a high electrical power density of 12 mW cm−2 at 180 °C was measured, among the highest values reported for cells assembled from crystalline solid electrolytes [75]. A phosphonate MOF that can operate over a wide temperature range was reported in 2010. Low-temperature proton conductivity was seen, even at −20 °C, which was facilitated with ordered, Zn-ligated water molecules. The phosphonate oxygen atoms assisted in anchoring and ordering water molecules into chains. Though proton conductivity was not particularly high (3.5 × 10−5 S cm−1 at 98% RH, 25 °C), a very low activation energy of 0.17  eV for proton transfer was measured through ordered interlayers [71]. Water-assisted proton conductivities of 1.45 × 10−2 S cm−1 at 363  K/95% RH and intrinsic, sustainable anhydrous proton conductivity of 1.1  ×  10−5  S  cm−1 at 503  K was measured for another zirconium phosphonate MOF. The protons can be transported with decent conductivity over a wide temperature range from 293 to 503 K [76]. Crystalline structure may change during transport of protons such to assist the conductivity in a variety of ways. In many cases, ions or molecules may be released from the structure with application of heat. These transitions are less frequently reported, but are nonetheless essential to gain an understanding of the proton transfer mechanism. Single crystal-to-single crystal (SCSC) transformations may be associated with water coordination. A Co-Ln MOF undergoes a crystalline phase transition at 25 °C and 95% RH, upon which a proton is released to increase the free charges and enhance the conductivity to reach 3.05  ×  10−4  S  cm−1, which is two orders of magnitude larger than that for the powder sample, demonstrating that the direction of hydrogen-bonded modes is the preferred proton conduction pathway in the crystal structure. A SCSC transformation was seen in an anionic Eu-based MOF, upon which a proton is transferred between the anionic form and the neutral version of the identical framework [51]. Here, phosphonate groups form parallel hydrogen-­ bonded chains, resulting in a proton conductivity of 1.25 × 10−3 S cm−1 at 150 °C, and water-assisted proton conductivity for a compacted pellet of micro-sized crystals attains 3.76 × 10−3 S cm−1 at 100 °C and 98% RH. Other unique attempts to determine the mechanism of transport in MOFs have been reported, such as the luminescent visualization of conductivity via shift in emission wavelength from a ruthenium (II) metallo-ligand [57]. A Zn-based MOF also showed SCSC transformations that were induced by loading of guests, resulting in structures that operated over a wide temperature range (−40 to 125 °C) [42]. In some cases, a nonporous network may be preferred to allow for more, closely spaced sites for ions to hop between. In 2013, it was shown that a MOF could be made such that it retained properties of porosity and the density inherent in solids to provide the elements of both to a useful proton conductor. The structure was formed around a bulky template, benzimidazole, and the authors demonstrate conductivity along the MOFs interior surfaces, rather than directly through the pore. Once water was removed and the crystalline shape changed, methanol was introduced, resulting in a 24× increase in conductivity due to accessibility of the methanol. This is in spite

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of methanol’s poor proton conductivity and was reasoned to have origins in the interaction between methanol and the framework phosphate groups [5]. Structural defects were utilized in a nonporous coordination polymer to incorporate mobile, uncoordinated H3PO4, H2PO4−, and H2O. EDX element mappings for Zn and P show homogeneous distributions of the elements, confirming the homogeneous distribution of H3PO4 throughout the defects. Embedding the phosphoric acid significantly lowers activation energy, as shown in Fig. 5 (right) [73]. Cheetham et al. found that a dense MOF will undergo an insulator-to-proton conductor transition upon hydration. Exposure to water results in a huge increase in conductivity at 17  °C to 3.9 × 10−5 S cm−1. The large associated activation energy (0.64 eV), as well as the small humidity dependence in the RH range 67–33%, suggests that protons do not transport through free water molecules adsorbed on the surface but through the framework and/or micropores. Furthermore, as the conductivity depends on humidity, the mobile ions are judged to be protons, not other ions such as Li+ [77]. There is some concern as to the role of the outer vs. inner surface of MOF crystals and the influence of grain boundaries between micro and nanosized crystals on the conductivity. Control of MOF morphology is an ongoing theme for many applications as well as increased understanding of the role of defects in MOF properties including conduction. To explore this, Cu MOF-based thin films were made. This format reduced the overall resistance of the electrolyte membrane and increased the efficiency of the device. Proton conductivities of the MOF nanofilm at RH 95% were measured, and the activation energy of proton conductivity was calculated to be 0.28  eV.  The highest recorded conductivity value for the nanofilm was 3.9 × 10−3 S cm−1 at 98% RH [45]. This demonstrated that the external surface of the crystals plays a measurable role. Shimizu also broached this topic. Synthesis of a series of Lanthanide MOFs, [Ln(H5L)(H2O)n](H2O) (L = 1,2,4,5-tetrakis(phosphon omethyl)benzene, Ln = La, Ce, Pr, Nd, Sm, Eu, Gd) resulted in MOFs with variable crystal sizes. The La and Pr complexes, which yield the biggest crystals within this series, are the best conductors of the group. This result contradicts previous concerns over degradation of MOFs at grain boundaries, which may result in small particles that enable proton conduction via extrinsic pathways. If this were the case, the ball milled samples would in fact exhibit the highest conductivity values [59]. Controlled hydrophilicity is also important as proton conductivity increases with hydrophilicity, but it is more difficult to retain high humidities during operation of a device. Authors report that a hydrophilic MOF, {NMe3(CH2COOH)} [FeCr(oxalate)3]·nH2O, strongly absorbed a large number of water molecules to result in high proton conductivities of ~10−4 S cm−1, even at a low RH of 65% [54]. Also, an anionic layered structure based on sulfonated indium, incorporating hydrogen-­bonded dimethylammonium cations and water molecules, was shown to exhibit conductivity over 10−3 S cm−1 at 25 °C and 40% RH, a very high proton conduction value for low humidity and moderate temperature along two axes of the crystals [56]. A Eu MOF exhibits temperature-dependent but humidity independent conductivity that is highest at 150  °C when associated water molecules are activated, and drops after loss of water. However, it can be rehydrated and reused [52].

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Although MOFs can be engineered with some degree of certainty, a prediction of the type and the packing of guest molecules in the coordination space of MOFs can often be difficult. A unique approach was taken by Nagarkar et  al. Since a large number of MOFs have been reported in the Cambridge Structural Database (CSD), the compounds within CSD were screened for the presence of high-boiling proton carriers, higher carrier concentrations, the extent of hydrogen bonding, inexpensive starting materials, and ease of synthesis. An oxalate-based type of MOF {[(Me2NH 2)3(SO4)]2[M2(oxalate)3]}n was found to be the most suitable candidate. Inspired by this, a new complex, {[(Me2NH2)3(SO4)]2[Zn2(oxalate)3]}n was synthesized solvothermally. The compound consists of an anionic framework [Zn2(oxalate)3]2−n that is interpenetrated with a cationic supramolecular net [(Me2NH2)3SO4]+n, which is formed by electrostatic and hydrogen-bonding interactions between sulfate anions and dimethyl ammonium cations. The proton conductivity was found to be 7  ×  10−5  S cm−1 at 30  °C under inert atmosphere, increasing to a maximum of 1.0 × 10−4 S cm−1 at 150 °C with an activation energy of 0.129 eV. Furthermore, the conductivity reached a maximum value of 4.2 × 10−2 S cm−1 at 98% RH [74].

4.2  Hydroxide Ion Conductivity Currently, alkaline fuel cells utilize the electrolytes KOH and NaOH for transport of hydroxide ions. The required chemical sources here are hydrogen as the fuel at the anode (just as in proton conducting membranes) and both oxygen gas and water at the cathode. The hydroxide is stripped from water at the cathode from an input of oxygen and water. At the anode, hydrogen gas combines with hydroxide to reform water (See Fig.  1). This was utilized successfully on the Apollo Mission, even before proton conducting membranes were used, despite the lower ionic mobility of the bulkier hydroxide ions. Since hydroxide ions in the electrolyte are compatible with less expensive, nonprecious metals (such as nickel electrodes versus platinum required in protonic systems), these would offer a significant economic advantage compared to proton conducting fuel cells. Furthermore, the kinetics of the reaction in alkaline fuel cells are faster and require less total catalyst. Polymer Nanocomposites There are fewer reports of hydroxide-conducting membranes, as organic polymeric systems that are chemically stable to the required conditions are rare, but interest in them remains. A review of hydroxide-conducting polymers in 2011 explains that there are three classes of polymeric materials which have been utilized as anion-­ exchange membranes in alkaline fuel cells [86]. These include polymer blends/ composites based on PEO-type materials or polybenzimidazole, organic–inorganic hybrid membranes (of interest here), and semi-interpenetrated networks. The following cations have been utilized as functional groups to bind to hydroxide in these

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anion-exchange membranes: ammonium, imidazolium, guanidinium, pyridinium, tertiary sulfonium, phosphonium, benzimidazolium, and pyrrolidinium. These classes of ions are all sensitive to OH− chemical attack, however. Fumapem (perfluorinated sulfonic acid/PTFE) is currently the go-to polymeric membrane used for this purpose as it demonstrates chemical stability similar to that of Teflon. The interested reader is directed toward a contemporary review that discusses perfluorinated sulfonic acid polymers and provides new insights [87]. A few polymer nanocomposites have been made and tested for hydroxide exchange. For example, cross-linked polyvinyl alcohol/ poly(diallyldimethylammonium) chloride was infiltrated with nanosized graphene oxide. The resulting composite exhibited both an increased mechanical stability and a hydroxide conductivity that increased by 80–120%, having values of 0.0121 S cm−1 @ 30 °C and 0.021 S cm−1 @ 80 °C [88]. Another example from 2017 is found in a pre-designed, hybrid core–shell nanoarchitecture, having nanoparticles composed of SiO2 cores and quaternary ammonium-functionalized polystyrene shells. The hydroxide-conducting groups are locally concentrated on the high surface area of the functionalized nanoparticle composite. Upon embedding these nanocomposites (20–70 wt %) into a polysulfone matrix, which is itself nonionic and resists aggregation, an extremely high hydroxide conductivity value was measured at 0.1881 S cm−1 at 80 °C. The well-connected ion channels that result are expected to be the reason for such high values. Additionally, favorable mechanical properties are observed, comparable to that of Nafion® [89]. MOFs Hydroxide ion conducting MOFs are rare. ZIFs are the prototypical MOFs used for hydroxide conduction, owing to their stability at high pH as described earlier. The first report in 2014 showed that the inclusion of alkylammonium hydroxides as ionic carriers into the pores was successful. Though the conductivity was not significantly high, an increase by four orders of magnitude (2.3 × 10−8 S cm−1 at 25 °C) vs. the unfilled ZIF-8 was shown. Hence, an otherwise insulating and neutral ZIF-8 became a hydroxide conductor. However, the Ea value of 0.70 eV is much higher than that of hydrated OH− ions in liquid (less than 0.2 eV), and suggests that there are some unfavorable features in the conducting pathways. For example, it was noted that the pore apertures are smaller than other MOFs at 3.4 Å in ZIF-8 [83]. Another report in 2014 showed that an aminated polymer could be threaded through ZIF-8. This was a slow process, as the monomer was incorporated by diffusion over a 1-month period before in situ polymerization was initiated, which was followed by in situ amination. Essentially, an anion-exchange membrane was formed within ZIF-8, which allowed for fast ion exchange under anhydrous conditions, although conductivity values were not reported [84]. In 2015, an ionic liquid (choline hydroxide) was incorporated into the pores of ZIFs and made into PVA composite membranes, which were studied for their hydroxide-conducting properties in low humidity conditions (33% RH). The ionic liquid is capable of maintaining hydration

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such that a membrane can operate at low humidity, which is also aided by the hydrophilic PVA. The ZIFs behaved simply as fillers here, and relatively high conductivities were measured at 33% RH at either 25 °C or 60 °C, with the highest conductivity measured as 9.8 × 10−4 S cm−1 at 60 °C [85]. In 2015, PSM with KOH was used to form defects in MOFs by replacing a linker with hydroxides, and these “missing linker defects” resulted in a four orders of magnitude increase due to the increased basicity and hydrophilicity of the resulting frameworks [81]. The conductivity values of these Ni-based MOFs increased up to 1.16 × 10−2 S cm−1, and the activation energy dropped down to 0.20 eV at 100% RH for the MOF modified with KOH. PSM was also utilized to form cationic MOFs. In 2014, Feng et al. showed that stripping a neutral framework of anions in a Cr-based MOF leads to cationic frameworks and mobile hydroxide ions [79]. Using this procedure, another cationic MOF preparation was reported. (Fe-MIL-101-NH2)+Cl− was shown to have excellent hydroxide conductivity and alkaline stability properties. An anion-exchange membrane was prepared from MOFs entrapped within porous bromomethylated poly(2,6-dimethyl-1,4-phenylene oxide) to yield a cationic MOF with a polyvinyl alcohol (PVA) coating. Upon soaking in NaOH, the entrapped cationic MOFs behave as the active OH− conductors, while the PVA coating is used to prevent crossover of fuels. A high hydroxide conductivity of 0.145 S cm−1 was measured at 353 K [80]. In 2016, Ghosh et al. showed the first MOF exhibiting conductivity via in situ incorporated hydroxides through the matrix, which were introduced as part of the initial synthesis. The highest hydroxide conductivity in a MOF was measured at 8 × 10−5 S cm−1 at 27 °C and 99% RH, and the activation energy was found to be 0.19 eV, which is reported to be similar to the transport of hydroxide ions in solution with no interference. The frameworks contained one-dimensional hexagonal channels extending through the structures, which were filled with hydroxide ions and water molecule to form an extended, hydrogen-bonded supramolecular chain of hydroxides and water molecules, which was theorized to be the reason for the OH− ion conductivity of the material [82].

5  Batteries Compared to fuel cells, batteries are currently the leaders in the energy storage sector. Batteries are considered to be of vital importance for portable energy solutions due to their high energy density and specific energy. In traditional aqueous batteries, such as 1.5 V alkaline or 2.0 V lead-acid batteries, the electrolyte is an aqueous solution of KOH or H2SO4, respectively. In primary batteries, the device only undergoes a discharge reaction, as the chemistry does not readily lend itself to operate in reverse. In secondary batteries, the cycle may be reversed for rechargeability and reuse. However, there is a limit to how many times a battery can be cycled as the chemistry at the electrodes is not infinitely reversible and, in general, the charge and discharge process degrades the electrodes over time. Furthermore, water breaks

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down at higher voltages, and therefore this limitation inhibits the energy storage capacity for aqueous batteries. Great interest exists in secondary batteries with high voltages to give lightweight devices, driven by advancements in cell phones, laptops, smartwatches, etc. Hence, the non-aqueous electrolytes with higher voltage stability that are necessary for lithium ion batteries are the focus of intense research efforts.

5.1  Lithium Ion Batteries Lithium is an ideal anode material for high energy batteries for two major reasons: it is the lightest elemental metal and it possesses a very large negative standard reduction potential (−3.0  V vs. NHE). This affords lithium batteries a superior energy storage capability compared to aqueous systems. Organic liquid electrolytes (e.g. ethers, carbonates) are often used as they offer stability at the operating potentials; however, they present a safety risk due to their potential flammability after a catastrophic event (e.g. a short circuit whereby significant energy could be released in a very short time period). Although invaluable for high energy primary batteries, lithium metal anodes do not recharge effectively, forming high surface area dendritic structures upon repeated cycling that are reactive with the electrolyte and that can, potentially, short the battery. However, the discovery of highly reversible Li+ intercalation into suitable anode (e.g. carbon) and cathode (e.g. metal oxides) materials has allowed for the development of Li-ion rechargeable chemistries that operate with similar energies to the primary batteries but avoid the formation of Li metal in the electrochemical reactions (see Fig. 6, top for a simplified schematic). Lithium ion batteries have been a boon for portable electronics. The use of ion intercalation electrodes is one of the most important advances made in the electronics industry, and the first major change in battery technology from the basic Galvanic cell model of metal ions reducing at the cathode and oxidizing at the anode. In most commercial devices, the electrolyte consists of lithium salts with a suitable anion to allow for solubility in the non-aqueous solvents used (e.g. LiPF6 or LiBF4 in ethylene carbonate/diethyl carbonate). Conductivities are relatively high in these liquid systems, approximately 10−2 S cm−1 at room temperature. However, the liquid electrolyte still presents a flammability risk so the search for solid-state electrolytes that are compatible with lithium ion batteries is well underway in an attempt to prevent leakage, flammability, and solvent breakdown at the electrodes. A typical figure-of-­ merit for a good solid-state Li+ conductor is one with conductivities that are above 10−4 S cm−1 at room temperature [6]. Although lithium ions offer higher energy densities than other alkali metal ions due to their favorable mass/voltage, sodium and potassium ion batteries are considered to be attractive alternatives due to the high abundance of these elements. Additionally, magnesium ion batteries are considered to be the “holy grail” for some [91], but there are very few options for liquid electrolytes and magnesium ions are thought to move too slowly in the solid state. Relatively minimal, but notable, work has been conducted on these non-lithium-containing systems and will be briefly discussed here.

Polymer Nanocomposites for Ion Transport

111 Charge

e-

e-

Discharge

Cathode

Aluminum

Discharge

Separator

Anode

Copper

Li+

Li+ Charge

MIT-20d

MIT-20-X

Ionic conductivity in MIT-20-X Fig. 6  Top: Schematic of an operating lithium ion battery. The anode and cathode are typically composed of graphite and metal oxides, respectively (Reproduced from Ref. [13] with permission from The Royal Society of Chemistry). Bottom: MIT-20d represents the neutral Cu2Cl2BTDD MOF, where bis(1H-1,2,3-triazolo[4,5-b],[4′,5′])dibenzo-[1,4]dioxin  =  (H2BTDD), with 1-D channels of 2.2 nm width. Upon treatment with LiCl, NaSCN, or MgBr2, the anion binds to the framework, leaving free cations to transport through the channels (Reprinted with permission from ref. [90])

5.2  Polymer Nanocomposites For Li-Ion Solid and gel polymeric electrolytes have been reported and some have found commercial use already. Significant progress has been made, and the reader is referred to reference [13] for a thorough review. The earliest examples are the lithium salts incorporated in PEO (PEO-LiX), which is classified as a dry solid polymer electrolyte (SPE). X is often chosen as a bulky group that results in relatively low lattice

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energy of the salt, such as BF4− or PF6−. Single-ion conductive polymers have been developed to allow for large transference numbers. Mixed conductor composites have found application in serving as advanced electrode materials, and in 2015, a review specifically describing the types of polymer nanocomposites utilized was published [92]. The use of inorganic fillers to improve the mechanical properties of polymers and inhibit PEO crystallization to allow for ion transport at room temperature has been applied yet again for lithium conductivity. In 1998, Croce, et al. demonstrated that inorganic nanofillers can be used as a type of “solid plasticizer” and allowed for room temperature conductivity. They utilized 13 nm TiO2 nanoparticles and 5.8 nm Al2O3, nanoparticles, resulting in large conductivity enhancements, ranging from 10−3 and 10−5 S cm−1 versus the range 10−4 to 10−8 S cm−1 for the ceramic-free electrolytes in the temperature range 30–80  °C.  They also measured an average Li+ transference number of ~0.6  in the 45–90  °C temperature range, which was the highest at the time for PEO-LiX [4]. Similarly, silica nanoparticle/PEO composites were described, whereby monodisperse 12 nm silica nanoparticles were grown in situ. Conductivities of 1.2 × 10−3 S cm−1 at 60 °C and 4.4 × 10−5 S cm−1 at 30 °C were reported. An electrochemical stability window up to 5.5 V was observed [93]. Ceramic particles are often considered to be inert fillers, yet their exact role remains under investigation. The first reports simply used this method to improve mechanical stability of the films. Wieczorek et  al. suggested that small particles (below 4  μm in size), allowed for enhanced interactions between the matrix and surface, owing to the significant increase in surface area [94]. This was followed by a study in 1998 describing the particle surfaces as binding sites for lithium ions, which behave as hard Lewis acids [95]. Lewis acid–base reactions on the surface of the particles are suggested to play an integral part in the conductivity. Further investigation by Croce, et al. demonstrated that the inorganic nanofillers behave as crosslinkers for the PEO and X-anions, stabilizing their separation and providing a modified matrix for lithium ion transfer [96]. More recently, solid composite polymer electrolyte with Y2O3-doped ZrO2 nanowires having positive-charged oxygen vacancies resulted in 2-order-of-magnitude increase in lithium ion conductivity [97]. In 2001, it was reported that ZnO nanoparticles (3.5 nm) having acetate groups on the surface cooperate with the PEO segments and lithium ions to form cross-­ linked structures, decreasing the film’s crystallinity and enhancing conductivity [98]. Nanosized fillers have an advantage over bulk materials in that they disperse more readily and reduce grain boundaries, while providing more free space for polymeric chains to move. Yet, they may precipitate from the matrix at high concentrations. In order to circumvent this problem, researchers designed a three-­ dimensional nanostructured hydrogel-derived Li0.35La0.55TiO3 framework, which was then used as a 3D nanofiller for PEO and displayed improved Li-ion conductivity up to 8.8 × 10−5 S cm−1 at room temperature [99]. The microstructure of the matrix is also of importance. A nanocomposite polymer electrolyte based on vinylidene difluoride-hexafluoropropylene copolymer

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(Viton) and TiO2 nanoparticles was formed into a lotus root-like porous structure, which resulted in high ionic conductivities of up to 1.21 × 10−3 S cm−1 at room temperature, with high electrochemical stability potential of 5.52 V (vs. Li/Li+), and a lithium ion transference number of 0.65 [100]. Carbon nanotubes have also been used for their tensile strength and high aspect ratios, as discussed previously. As they are not electrically insulating, they were packaged within insulating clay layers to form effective 3D nanofillers, increasing the lithium ion conductivity of a PEO electrolyte by almost two orders of magnitude, while improving the mechanical properties significantly, even with a very modest 5 wt % loading [101]. Other Ions Lithium is of limited availability, and the global distribution of lithium is un-even, appearing in the highest amounts in South America. Sodium ions are considered a cost-effective alternative as there are no imminent shortages of it [102]. Sodium-­ based batteries may be thought of as complementary to lithium ion batteries, rather than competitive with them as the specific energy is lower than that of lithium ion batteries, limiting the number of applications. However, in stationary settings, such as power back-ups in factories or power plants where this is less of an issue, they could provide a significant, cost-effective alternative. The insertion chemistry of sodium is similar to lithium, but the electrodes need to be optimized and new electrolytes must be found. To date, only a few solid-state electrolytes have been reported. In 2010, Kumar and Hashmi reported on a novel sodium ion conducting, gel polymer electrolyte nanocomposites based on poly (methyl methacrylate) (PMMA) with dispersed silica nanoparticles at ~4 wt. % [103]. The authors report that these gel electrolytes have a maximum conductivity of ~3.4 × 10−3 S cm−1 at 20 °C, while providing both the expected mechanical improvements, and thermal and electrochemical stability. Similar to PEO-LiX, a PEO-NaPO3 has been described, and subsequently loaded with ceramic BaTiO3 nanofillers, exhibiting a two-order of magnitude enhancement vs. the nanofiller-free film, reaching a maximum ionic conductivity of 1.2  ×  10−6  S  cm−1 at 345  K with a  cationic transport number of 0.33 [104]. Most recently, magnesium ion batteries are gaining attention in the secondary battery arena. With the relatively high abundance of magnesium vs. lithium, the high energy density of magnesium, and the comparable size of the two ions, magnesium batteries would be highly valued. Additionally, magnesium metal is less reactive than lithium in a battery. The major obstacle as of now is in developing a suitable magnesium ion conducting electrolyte. All liquid electrolytes haven proven too corrosive to the electrodes, and focus has gone directly to the stage of developing solid-state electrolytes. The latest successes have mostly been with all-inorganic frameworks, such as magnesium scandium selenide spinel structures [91]. However, reports that utilize similar recipes to the PEO-LiX systems discussed above by incorporating ceramic nanoparticles are in the literature. For example, in 2017, a composite polymer electrolyte was reported [105]. The authors demonstrated that a

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mixture of PEO, Mg(ClO4)2, nanometer-size silica fillers, and an ionic liquid (1-Ethyl-3-methylimidazolium bis(fluorosulfonyl)imide) shows high ionic conductivity of 5.4 × 10−4 S cm-1 at ambient temperature and exhibits a wide voltage stability range of 4 V vs. Mg/Mg2+.

5.3  MOFs MOFs are naturally finding use in batteries and supercapacitors for similar reasons as in proton conducting membranes [106]. MOFs were first employed in battery technology as electrode materials. Notably, MOFs were used as porous metal organic sources that could serve as sacrificial templates to yield porous metal oxides for lithium ion intercalation [107]. Only recently have MOFs found use as lithium ion electrolyte materials. In 2011, the first example was described by Long et al. Here, it was demonstrated that the incorporation of a typical electrolyte solution (containing lithium tetrafluoroborate in ethylene carbonate and diethyl carbonate solvent) into a MOF with open metal cation sites can produce a solid with significant ionic conductivity. Furthermore, it was found that first anchoring lithium isopropoxide to the unsaturated magnesium ion sites in the MOF was necessary to yield a higher concentration of lithium ions. Specifically, after soaking the original MOF, Mg2(dobdc), (dobdc = 1,4-dioxido-2,5-benzenedicarboxylate) in the lithium isopropoxide, then filling with lithium tetrafluoroborate, a pellet of this composite material was found to have a conductivity of 3.1 × 10−4 cm-​1 at 300 K, with an activation energy of just 0.15 eV [108]. In an interesting extension to this, the Long group also demonstrated the first solid-state, room temperature magnesium ion conductor [109]. A series of solid magnesium electrolytes were synthesized with the use of magnesium phenolates to initially bind with coordinatively unsaturated metal sites lining the pores of the two MOFs: Mg2(2,5-dioxidobenzene-1,4-dicarboxylate) and Mg2(4,4′-dioxidobiphenyl-3,3′-dicarboxylate). Room temperature ionic conductivities up to 2.5 × 10−4 S cm−​1 were measured. As ionic channels are essential features in biological cell membranes, selectively transporting ions, some researchers were inspired by this idea. Similar to Hupp’s report of HKUST-1 for proton transport, wherein open, acidic copper sites result in an increase in acidity of coordinated waters [10], Shen et al. demonstrated Li+ transport in HKUST-1. Here, LiClO4 in propylene carbonate was introduced into the MOFs, whereby ClO4− ions spontaneously bind to the open metal sites, forming negatively charged MOF channels and enabling transport of Li+ ions with a low activation energy of 0.21  eV.  They also demonstrated this approach successfully with Zr-containing MOFs having open metal sites, such as UiO-67 (Zr(IV) biphenyl dicarboxylate MOF). The authors measured a low activation energy of 0.12 eV for lithium ion conductivity, which reached ~10−3 S cm−1 with a lithium transference number of 0.65. Notably, when the same procedure was utilized to introduce Li+ into the non-coordinating IRMOF-1, having no open metal sites, activation energies were significantly higher (0.4–0.5 eV), demonstrating the active role taken by the

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framework in the conductivity mechanism. The authors also note that smaller pore sizes in the MOFs tend to improve conductivity [28]. A recent article demonstrates the first use of a MOF that functions as a matrix to hold ionic liquids for lithium ion transport. Kitagawa et al. showed that ZIF-8 could be doped with a mixture of 1-ethyl-3-methylimidazolium bis (trifluoromethylsulfonyl) imideand LiTFSI (lithium bis (trifluoromethylsulfonyl) imide). It was suggested that the Li+ cations diffuse through the micropores via the exchange of the solvating counter anions, which is comparable to the Grotthuss mechanism used to describe proton conductivity [110]. This was followed by another report in 2017, whereby an ionic-liquid-impregnated MOF demonstrated a high room temperature ionic conductivity of 3.0 × 10−4 S cm−1, a Li+ transference number of 0.36, and good compatibilities against both Li metal and active electrodes with low interfacial resistances [111]. In related work involving doping a MOF with ion-conductive guests, Chen et  al. reported a polystyrene sulfonate and lithium salt could be threaded through a HKUST-1 membrane for fast and selective lithium ion separation (Li+ vs. Na+, K+, and Mg2+) [112]. The resulting composite exhibited very high Li+ conductivity of 5.53 × 10−4 S cm−1 at 25 °C, and 1.89 × 10−3 S cm−1 at 70 °C, which are five orders of magnitude higher than that of the pristine HKUST-1 membrane. The selectivity was attributed to the different size-sieving effects and the affinity differences of the Li+, Na+, K+, and Mg2+ ions to the sulfonate groups. In 2017, A Cu(II)–azolate MOF that contains tubular pores was shown to support record-high single-ion conductivities for Li+-, Na+-, and Mg2+-loaded MOFs. The framework was shown to undergo a reversible single crystal-to-single crystal (SCSC) transition between neutral and anionic phases upon reaction with stoichiometric amounts of salts. The stoichiometric transformation between the two phases allows loading of record amounts of charge-balancing Li+, Na+, and Mg2+ ions for MOFs. Whereas the counter anions are bound to the metal centers and thus stationary, the cations move freely within the one-dimensional pores, giving rise to single-­ ion solid electrolytes. The respective Li+-, Na+-, and Mg2+-loaded materials exhibit ionic conductivity values of 4.4  ×  10−5, 1.8  ×  10−5, and 8.8  ×  10−7  S  cm−1. With addition of LiBF4, the Li+ conductivity improves further to 4.8 × 10−4 S cm−1. These are the highest values yet observed for MOF solid electrolytes (See Fig.  6, bottom) [90]. Novel scandium-based alkaline MOFs were prepared and tested for ionic conductivity. The newly reported structures, {[ScM(μ4-pmdc)2(H2O)2]·solv}n, where M = Li, Na; pmdc = pyrimidine-4,6-dicarboxylate; solv = corresponding solvent, compounds were made such that they contained ions in the matrix. However, the conductivity was low in the pristine MOFs, but it was further enhanced by doping with divalent transition metal ions and soaking the samples in alkaline salts, resulting in an increase in the number of alkaline ions. This resulted in high conductivities for MOFs at room temperature, measuring 4.2  ×  10−4 and 9.2  ×  10−5  S  cm−1 for M = Li and M = Na, respectively [113].

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6  Supercapacitors Supercapacitors are another type of energy storage material, and are often grouped with batteries when describing the most promising applications for modern, green energy storage solutions. Batteries are known to have a high specific energy, but low power density. In contrast, supercapacitors provide a lower energy density but much higher power density. Although the two are often posed as competitors, batteries, and supercapacitors but may also be coupled to obtain the properties of both [114]. An important and desirable aspect of supercapacitors is they can be charged and discharged rapidly, as they use electrostatics to store charge, rather than internal reactions of ions between electrodes, to obtain an immediately available energy supply. Furthermore, as there is no chemical reaction to consider, they are theoretically nearly infinitely cycleable with little degradation. An example lies in the use of all-­ electric hybrid buses. The supercapacitors can be used to quickly accelerate the buses and recharge upon braking (which happens frequently on bus routes), but do not have enough energy to power the bus alone. Hence, batteries provide energy storage for the bulk of the ride. Also, buses and vehicles that are powered only by supercapacitors, called “capa” vehicles, are also possible. These contain supercapacitors under the seats, can be partially powered by braking and while stopped at bus stops via power collectors, and are less expensive than lithium ion batteries, therefore holding great promise for use in situations where a route is known and stops are planned [115]. In the basic design of any capacitor, two conductors are separated by a dielectric that can be polarized in the presence of an electric field. In a simple dielectric capacitor, metal plates are separated by a dielectric. When charged, the positive charge moves to one of the metal plates, while the negative charges move to the other, creating an electric field that results in the alignment of dipoles in the dielectric. The measured capacitance for typical solid-state electronics is on the order of picofarads to microfarads, limiting the functionality of these for energy storage. Since the capacitance is an extrinsic property that depends on the amount of material, higher surface area would yield a higher concentration of charge, allowing for smaller devices. Supercapacitors operate using the adsorption/desorption of ions in an electrolyte that is in contact with much higher surface area electrodes (i.e. activated carbon), but on a grander scale, resulting in capacitance values on the order of Farads, which is one million or billion times larger than standard dielectric capacitors. Yet again, nanotechnology has been instrumental in advancing the utility of capacitors, as the nano-featured surfaces provide the requisite high surface area. The individual, metal-coated porous plates are soaked in an electrolyte and placed next to one another with a thin insulator between. Upon charging, the two sides develop a charge separation, and the electrolytes polarize with the charges aligned on either side. Because both dielectric layers may be as thin as one molecular layer, such electrolytes are often referred to as double-layer dielectrics. The electrolytes used will change the operating characteristics of the supercapacitor, including the voltage and capacitance. Due to the high concentration of charges, aqueous solutions such as of sulfuric acid have traditionally been used in the past, but the aqueous phase limits the voltage possibilities.

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6.1  Polymer Nanocomposites Neat polymer dielectrics can be used, the most common of which are poly (4-­vinylphenol) (PVP) and poly(vinylidene) fluoride (PVDF), as they offer an advantage in high breakdown strength when compared to other materials used for capacitor applications. However, the low dielectric constants significantly limit their application potential. Perovskite ceramic nanoparticles offer higher dielectric constants and can be combined with polymers to yield improved polymer nanocomposites [116]. In order to increase the processability and effective surface areas, BaTiO3 nanoparticles, and modifications thereof, such as Ba0.5Sr0.5TiO3, have been functionalized with PVP to result in a more homogenous product having an increased dielectric permittivity of 77 at 1 kHz, 28% higher than when measured with untreated BaTiO3. It was shown that modification of the surface is also important for close contact between the polymer dielectric and ceramic. Here, hydroxyl groups on the surface of BaTiO3 nanoparticles served to hydrogen bond with the fluorine atoms in the polymeric PVDF [117]. Another method to maintain close contact of the particle and matrix is to graft the matrix material directly onto the nanoparticles. For example, PVDF copolymer chains were grown off of TiO2 nanoparticles. These particles could then be easily dispersed into a polymer matrix and cast into thin films, which resulted in doubling of the energy storage capabilities (from 3.25 to ~7.0 J/cm3) with 50% loading [118].

6.2  MOFs MOFs present extremely high surface areas and lend themselves well to this application. Some work has been carried out with decomposed or pyrolyzed MOFs to yield nanoporous metal oxide structures or porous carbon materials that are comparable to activated charcoal [119]. As for previous devices mentioned throughout the chapter, the framework structure itself can be used as-prepared or modified with PSM or guests to behave as the electrolyte. Since MOFs themselves typically exhibit low conductivity, (which results in an increase in the bulk electric resistance of such devices) PSM or guest incorporation has been utilized most frequently to address this. Various strategies have been taken to improve MOF–conductor contact, such as forming nanocrystalline-sized MOFs for increased interaction with the conductor, or interweaving the conductor through the MOFs. In 2014, Yaghi et  al. made nanocrystalline-­sized MOFs (termed nMOFs), which were doped with graphene. They tested various metal nodes and organic linkers and found that one particular system, Zr6O4-(OH)4(BPYDC)6 (BPYDC = 2,20 -bipyridine-5,50 - dicarboxylate, termed nMOF-867), exhibits an exceptionally high areal capacitance of 5.09 mF cm−2. The authors note that this is about six times that of the supercapacitors made from the benchmark commercial activated carbon materials. Furthermore, performance that is preserved for at least 10,000 charge/discharge cycles

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Fig. 7  Top: ZIF-67 as utilized for supercapacitor function. (a) Electron cannot migrate along or access the skeleton of MOF crystals, while electrolyte can get in and out of the MOF pores by diffusion effect. (b) After interweaving MOF crystals with conductive PANI, both electron and electrolyte can access MOF surfaces and an electrochemical double-layer is formed on the surface of PANI-ZIF-67-CC (Reprinted with permission from ref. [120]). Bottom: Construct for nMOF Supercapacitors (Reprinted with permission from ref. [121])

(compared to 5000 for batteries) was shown. See Fig. 7, bottom [121]. A flexible, solid-state supercapacitor utilizing cobalt-based ZIF-67 crystals (the linker is 2-­methylimidazole) that support the movement of the electrolyte through the pores has been reported in order to reduce the bulk electric resistance of MOFs. Polyaniline (PANI) chains were interwoven into the MOF crystals via electrochemical deposition, which were deposited onto carbon cloth (CC) before further electrically depositing PANI to give a flexible, conductive, and porous electrode (noted as PANI-ZIF-67-CC). This approach allows significant contact and electron transfer between the electrolyte (KCl) and the polymer. The underlying structure of the

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MOF remained intact throughout this process. The resulting PANI-ZIF-67-CC exhibits an extraordinary areal capacitance of 2146 mF cm−2 at a sweep rate of 10 mV s−1, and further, a functioning supercapacitor device was successfully made. See Fig. 7, top [120].

7  Solar Cells The first solid-state solar cells were developed in 1954 at Bell Labs and quickly found application, particularly in satellites. The devices soon advanced and the efficiency of conversion of photons to electricity increased from an initial 6% to 25% for these first generation silicon solar cells. Although they are based on an abundant element, it is required that high quality, monocrystalline silicon wafers are formed for high efficiency, making them costly to produce. Additionally, there are environmental concerns with their manufacturing process. In SiO2 the silicon must be reduced and treated with caustic chemicals at high temperatures for asymmetric doping with negative (i.e. nitrogen) and positive (i.e. boron) additives to create a charge gradient and allow electrons to travel directionally. Furthermore, flat, brittle sheets of silicon are inflexible, limiting their utility. Second generation cells consist of amorphous silicon and semiconductor compounds such as GaAs, and have reached similar or higher efficiencies (~10% up to 30%). The cost is lowered, but their production also requires significant energy consumption and is based on less abundant chemicals in some cases. The third generation of solar cell design includes dye-sensitized solar cells (DSSCs), first demonstrated by Grätzel in 1991 [122], and related organic and quantum dot-containing cells [123]. Intense research into this area is ongoing. A review article in 2015 pointed out that upon searching under the keywords “dye-sensitized solar cell,” more than five articles are published a day [124]. DSSCs are of interest because they are amenable to variations in manufacturing and can be adapted to a wide variety of device geometries. They also might require less intense manufacturing and not be constrained by demanding high purities compared to silicon. However, the organic components in current generation solar cells are not nearly as stable or efficient (maximum ~13%) [125] as silicon solar cells, but they continue to be investigated for their other favorable attributes. These solar cells have many parts to them, as the energy source is direct from the “tap”: A DSSC must absorb light, transfer electrons, support this with an electrolyte, and utilize a redox couple to replenish the organic molecule with electrons for continued use. One can rationalize that this complexity is relatively simple compared to photosynthesis and formation of fossil fuels, however. Essentially, photosynthesis, followed by the formation of fossil fuels over billions of years to yield hydrocarbons that power internal combustion engines is bypassed in DSSCs. See Fig. 8 for a representative schematic of DSSCs. The first DSSCs utilized a redox couple in a liquid electrolyte, sandwiched between electrodes to support the movement of electrons and replenish electrons in the organic dye, where I−/I3− is the classic couple system [124]. This is because

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eFig. 8  Schematic of a dye-sensitized solar cell (DSSC). First, light passes through transparent, conductive FTO glass and induces excitation of the dye molecules. The excited electrons are injected into the conduction band of the TiO2 semiconductor, leaving holes in the dye. The dye is regenerated upon electron donation from the reduced state of the redox couple (I−) and holes are transported through the electrolyte layer. Finally, the electrolyte is regenerated upon accepting electrons from the counter electrode. Reproduced from Ref. [126] with permission from The Royal Society of Chemistry

when the electron is transferred, a hole is left behind in the dye. There is a limitation to the stability and lifetime of these DSSCs, however, as the iodine couple is corrosive to the electrodes and has a low breakdown voltage of 0.9 V. Additionally, iodine absorbs light in the visible region, preventing this light from reaching the dye and reducing efficiency. It is essential that the dye is absorbing and can effectively transport electrons to the wide band-gap semiconductor, typically TiO2. It is also important for recombination of holes and electrons at the anode to be prevented, and the electrolyte also plays a role here. Much of the work on DSSCs has been focused on improvements to the dye sensitizer. However, stability and longevity of these devices hinge on the redox electrolyte’s properties. Such electrolytes are going through the same evolution process as other devices before: first liquids, then ionic liquids, and now the search for solid electrolytes. Recently, a new sub-generation of quasi- solid-state DSSCs utilizing the Co(II)/ Co(III) redox couple have renewed interest in this area as they open up the voltage-­ operation window and improve the lifetime of the electrodes [124, 127]. This newer redox couple, when coupled with ionic liquids, has helped in improving the performance of DSSCs, but ionic liquids still remain susceptible to leaks, especially when exposed to high temperatures under direct sunlight. Solid-state electrolytes, especially polymeric hole transporters, have been employed, but face issues with crystallization. Again, polymer nanocomposites and MOFs show improvement in the

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films mechanical strength and heat stability. Such nanocomposites are just one type of solid electrolyte under development. Perovskites and copper-based compounds, such as CuI, which are hole-transporting inorganics, are also showing promise. However, contact between the electrolyte and sensitizer is more difficult to achieve in the solid state. A review in 2016 discusses the need for a solid-state electrolyte, which is key to increasing the stability and lifetime of DSSCs [127].

7.1  Polymer Nanocomposites Nanostructured surfaces are useful in order to maintain significant contact between the electrolyte and electrodes, and to prevent electron–hole recombination. Many examples use the archetypal ion conducting polymer, PEO, and embedded ceramic composites to prevent crystallization of the matrix. For example, composites of PEO and titania nanocrystals resulted in an energy conversion efficiency of 4.2% at 65.6 mW cm−2, one of the highest values reported at the time (2002). This relatively high efficiency was attributed to the amorphous nature of the polymer and the increased contact of the polymer electrolyte with the dye sensitizer [128]. In 2004, a modification was made to enhance interfacial contact. The polymer chosen was poly(ethylene oxide dimethyl ether), or PEODME, which penetrates well into the titania/dye layer. Here, silica nanoparticles were used to prevent crystallization of the polymer, and the overall system, a composite consisting of PEODME/MI (M = K+, imidazolium+) / I2 / fumed silica nanoparticles, yielded favorable mechanical strength and a solar conversion efficiency of 4.5% at 100  mW  cm−2 [129]. Researchers also demonstrated the use of a hybrid PEO/PVDF/TiO2 nanoparticle system as solid-state electrolytes in DSSCs [130]. The iodide/triiodide redox couple was incorporated into poly(ethylene oxide)/polyaniline (PEO/PANI) solid-state electrolytes, aiming at expanding the catalytic event of I3− reduction from the electrolyte/counter electrode interface to both the interface and electrolyte system and shortening the charge diffusion path length. PANi is also responsible for dye regeneration and hole transfer to the counter electrode. The constructed DSSC with the iodine redox couple incorporated into PEO/1.0 wt% PANi electrolyte yields a maximum efficiency of 6.1% in comparison with 0.8% obtained from a PANi-free electrolyte-­based solar cell and 0.1% for a PANi-based solar cell [131]. This interest in polymer nanocomposites continues today. Armel et al. demonstrated that molecular plastic crystal (succinonitrile)-based solid-state electrolytes can perform very well in solid-state DSSCs, when combined with a porphyrin sensitizer and inclusion of nanoparticulate SiO2 [132]. This produced the highest device efficiency of 5.3% at moderate light intensity. In 2018, a few more publications have demonstrated the application of polymer nanocomposites for DSSCs. These include a polymer fibrous membrane electrolyte, whereby an electrospun nanocomposite polymer (PVDF-HFP + 6 wt% TiO2 nanofibers fillers) with Li+ exhibited a conductivity of 1.87 × 10−2 S cm−1 at room temperature. However, the function within solar cells was only suggested, not measured directly [133]. Additionally, PANI-grafted

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silica nanocomposite-based gel electrolytes for quasi-solid-state dye-sensitized DSSCs were made, where the silica was again used to prevent crystallization. An efficiency of 7.15% was measured [134].

7.2  MOFs Nanocrystalline Solar Cells. MOFs have only found limited use thus far in DSSCs. Copper-based MOFs have been employed as the sensitizer in two examples. In the first, HKUST-1 crystals were doped with iodine to improve the charge transfer across the surfaces of the DSSC, resulting in a low, but measurable efficiency of 0.26% [135]. Also, a similar solar cell was constructed with HKUST-1, whereby carbon nanotubes were incorporated with the titania layer for enhanced conversion efficiency, albeit a similarly low value. With MOFs only, the efficiency was measured at 0.20%, whereas incorporation of nanotubes increased this value to 0.48% [136]. In another example, a pillared Zn-porphyrin MOF was utilized to study the role of the MOF in a DSSC. Evidence was shown that clarified the MOF structure is indeed the active sensitizer in these DSSCs, as care was taken to remove any precursors that may remain in the MOF pores post-synthesis [137]. ZIF-8 has been grown on the surface of titania nanocrystals as an intermediate layer before coating with a dye. The resulting increased open-circuit voltage was ascribed to the inhibition of interfacial charge recombination by ZIF-8 [108]. Also, MOFs have again been used as decomposition templates to derive metal oxide anode materials with large surface areas for more surface contact with the sensitizer and electrolyte [138]. Bella et al. reported the first MOF-based polymer nanocomposite for electrolytes in DSSCs [139]. Here, a polymer composite was formed, containing a Mg-based MOF (Mg3(benzenetricarboxylate)2) mixed with PEG(dicarboxylic acid): PEG(maleic acid), which was activated by soaking in ionic liquids and utilized as an electrolyte for quasi-solid solar cells. An efficiency of 4.8% was reported, along with outstanding long-term durability. The MOFs were shown to interact with the TiO2 surface, effectively shielding trap states (electrons get stuck, rather than transferred) of the photoelectrode, resulting in higher photovoltage.

8  Conclusions and Outlook The scientific community is getting ever closer to developing safer, sustainable, high-performing compact devices using new, designer materials. All-organic polymers have provided optimism in the construction of solid state (or quasi-solid-state) devices, having some of the essential features that are in-demand for modern electrochemical solutions. Yet, the limitations were too great to achieve the robustness that is needed for many applications. Polymer nanocomposites have allowed for the

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promise of advanced solutions. As the improvements in computing and electronics continue, the development of compatible materials is attempting to keep up in this cooperative relationship. The movement of ions is at the crux of the technology’s growth and the corresponding electrolytes have reached the third stage of the maturation process—the solid stage. Polymer and nanomaterial designs are expanding and now that nanomaterials can be made with advanced understanding and control, they can be tailored to meet the needs of a specific device. With the arrival of MOFs and the explosion in synthetic studies and methodologies available, the scientific community now has modular platforms to experiment with. MOFs can be considered as a type of molecular Lego™, wherein functionalities can be included as desired with the choice of metals, linkers, shapes, sizes, and porosity, and are proving to be more than just a curiosity, but rather functional candidates for modern electronic components. Furthermore, the polymeric nature of MOFs, which are ordered to provide well-characterized pores, but contain easily-accessed defect sites that often continue throughout the matrix in low-density frameworks, provide a guide to study and develop the mechanisms of ion transport to enable development of even better materials in the future.

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Efficient Light Harvesting in the Nanotextured Thin Film Solar Cells Mohammad Mahdi Tavakoli

Abstract  Nanostructures have a great role for boosting the efficiency of the photovoltaic devices by trapping the light beams and increasing the absorption as well as carrier collection. This chapter summarizes the most recent light management approaches in the thin film solar cells. First, the effects of various nanostructures in terms of geometry, pitch size, and aspect ratio on the optoelectronics properties of devices are discussed in detail. Later, the applications of the nanostructures as anti-­ reflection layers and device substrates in the solar cells are explained. Then, their advantages and possible mechanisms are discussed to have a better understanding from the nanostructured-based devices. Additionally, the issues induced by nanostructures and the possible solutions for addressing these shortenings are proposed in this chapter. Keywords  Thin film solar cell · Light management · Nanostructure · Geometry · Efficiency · Light absorption · Anti-reflection · Self-cleaning property · Superhydrophobicity · Flexibility · Lambertian limit · Recombination · Carrier collection

1  Introduction Light management in a solar cell device is a key strategy to enhance the solar to electric power conversion efficiency (PCE). Integrating of solar cells with nanostructures has been focus of many research groups to revolutionize the device architectures. Nanotextured structures not only improve the light absorption properties but also reduce the thickness of light absorbers and materials consumption. Also, they can improve the charge carrier collection by controlling the mean free path of the carriers. Basically, the nanotextured patterns are associated with maximizing the short-circuit current density (Jsc) of devices. Depending on the device architecture, it can influence on other photovoltaic (PV) parameters such as open circuit M. M. Tavakoli (*) Department of Electrical Engineering and Computer Science, Massachusetts Institute of Technology, Cambridge, MA, USA e-mail: [email protected] © National Technology & Engineering Solutions of Sandia, LLC 2021 M. Alston, T. N. Lambert (eds.), Energy-Sustainable Advanced Materials, https://doi.org/10.1007/978-3-030-57492-5_5

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photovoltage (Voc). Extensive efforts have been exerted to improve the PV parameters using nanotextured substrates for light management purpose. In this regard, nanostructures can be stacked to the transparent side of the device as an anti-reflection layer or they can be employed as substrates for the device fabrication. In both cases, optimization of the geometry for these nanostructures such as periodicity, shape, and aspect ratio is required to efficiently improve the PCE of the devices. Without this step, the performance of device is compromised by large surface area inducing defects and additional charge recombination, particularly for device fabricated on nanotextured substrates. Consequently, fabrication of a uniform thin film on nanotextured structures can address some issues in this field such as the parasitic absorption and surface recombination. An overall view of nanotextured based solar cells will help to have a better understanding from the advantages of these architectures on the fabrication of high-performance devices. This chapter begins with the fabrication of nanostructures with optimum geometries for light management purpose. Then, their applications in the perovskite solar cells such as anti-reflection layer and device substrates have been explained.

2  Design of Nanostructures In order to fabricate efficient solar cell devices, the enhancement of light absorption is an important step, which needs to be taken into the consideration. In the past years, extensive studies on device architecture innovation have been performed to boost the device performance. For instance, anti-reflection layers, back-reflector coatings, and nanotextured substrates were investigated to minimize the reflection and trap the sun light within the device [1–3]. Among these techniques, utilization of nanostructures as a substrate or anti-reflection layer is an efficient approach to enhance the light absorption and device performance. For this purpose, there are various types of nanostructures such as nanocone (NC), nanopillar (NPL) or nanowire (NW), nanospike (NSP), nanowells, etc. These nanostructures are strong candidates for enhancement of light harvesting efficiency for  a solar cell [4–7]. The material consumption and the thickness of light absorber can be significantly reduced due to light trapping properties. Because nanostructures can improve the optical absorption and enhance the charge carrier collection [2]. To further improve the light harvesting properties of nanostructures, the proper design and optimization of geometry are key factors. On the other hand, fabrication of the nanotextured device may increase the complexity of fabrication process and it is a bit challenging as compared to planar structure. In some cases, the utilization of nanostructures increases the surface recombination due the increase in surface area, which can reduce the device performance. In order to design an efficient solar cell device, it is important to understand how much of the absorbed light comes from nanotextured active layer as opposed to the passive materials in the device such as electrodes, since the photo-generated carriers in active materials can only contribute into the photocurrent [8, 9]. Thus, optical absorption enhancement by nanotextured

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structures is an effective way to reduce parasitic absorption in the passive layers of devices [10, 11]. To fabricate these nanostructures, many different techniques were explored. They can be categorized as top-down, bottom-up, and their combination. Among these approaches, top-down method has the highest accuracy and controllability of nanostructure geometries (aspect ratio and ordering) [12]. Figure 1a–d show the schematic of NP, double-NP, multiple-NP, and NC arrays of germanium (Ge). As seen, the geometry of nanostructures can be tuned by applying various parameters during the growth step. It is noteworthy that the reflectance

Fig. 1  Schematics of Germanium nanostructure arrays with different shapes, (a) Single-NP array, (b) Double-NP array, multiple-NP array, and nanocone (NC). (e) Broadband-integrated absorption of 1000  nm pitch Ge multiple-NP array versus number of pillar. Red-dashed line indicates the broadband-integrated absorption of 1000 nm pitch Ge NC array. (f) Broadband-integrated absorption of Ge NC array at 1000 nm pitch [2]

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of each array simply depends on the filling ratio (FR) of the material, i.e., FR = πD2/3.4A2, where D and A are the diameter and the pitch of single-NP array, respectively. This relates the reflectance of photons to the top surface of NP array and thus the total reflectance significantly depends on the top surface area fraction or FR of NP array with similar diameters [13]. As a result, NC array has the lowest reflectance due to the minimum surface area. Figure 1e shows the optical absorption of NC array with multiple-NP structure with the same thickness (2000 nm), 1000 nm pitch size, and different segment (N). The broadband-integrated absorption and the ideal short-circuit current density (Jsc) of these arrays show that the maximum absorption and Jsc are obtained by increasing the segment number (N). As a result, NC array shows the highest values for both absorption (95.1%) and Jsc (58.4 mA/ cm2) as shown in Fig. 1e, due to the lowest surface area and reflection. Figure 1f illustrates the broadband-integrated absorption of 2000 nm height Ge NC array [2]. In this plot, the y axis is the bottom diameter of NC array and the x axis stands for the pitch size. It is clear that the reflectance of NC array is marginal regardless of bottom diameter and pitch size, since the cone tip diameter is assumed to be zero. Consequently, the largest bottom diameter for each pitch size shows the best optical absorption [14, 15]. Figure 2a depicts the schematics of NSP array before and after amorphous silicon (a-Si) solar cell fabrication. The angular-view SEM images of NSP substrates with 1.2  μm pitch size are shown in Fig.  2b1, c1, d1. The substrate is anodized aluminum oxide (AAO) fabricated through several anodization and etching steps. Notable, the pitch size can be controlled by applying different voltages. The geometry of nanostructures can be controlled by etching time and the concentration of solutions. Figure 2b2, c2, d2 illustrate the corresponding a-Si devices fabricated on NSP substrates with different heights controlled by anodization time (30 min, 3 h, and 6 h, respectively). As seen, there is structural transportation from nanoconcave to nanospike by increasing the anodization time. Due to their peak to valley height differences in Fig. 2b1, c2, d1, they are named as NC 200, NSP 600, and NSP 1200, respectively [16]. The optical absorption and photo-carrier collection are key factors for improvement of device parameters such as short-circuit current density (Jsc), open circuit photovoltage, (Voc), and fill factor (FF), which are coupled strongly. Consequently, investigation on the coupled optical and electrical properties of solar cell device based on nanostructures is of critical importance [17, 18]. Figure 2e, f show a systematic study on the geometry optimization of nanostructures and its effect on the optical and electrical properties of solar cell devices using finite difference time domain (FDTD). The integrated optical absorption of a-Si devices versus pitch size and height of NSP is shown in Fig. 2e. As it can be observed, the optical absorption of solar cell based on NSP array is much higher than that of the planar device. It is clear that by increasing the height of NSP array, the optical absorption is monotonically enhanced, where the highest absorption is 88.8% for 1.2 μm pitch size with 1.2 μm height. This indicates the higher light trapping effect for the higher aspect ratio. Figure 2f shows the integrated absorption in a-Si layer calculated by FDTD simulation, represents

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Fig. 2 (a) Schematics of nanospike (NS) array and multiple thin film layers deposited on NSP for solar cell fabrication. Angular-view SEM images of (b1) nanoconcave (NC 200), (c1) NSP 600, (d1) NSP 1200, and their corresponding a-Si solar cell at (b2), (c2), and (d2). (e) The integrated optical absorption under AM 1.5  G solar spectrum with different NSP pitches and heights. (b) Integrated optical absorption and its calculated Jsc under simulated AM 1.5 G solar spectrum in a-Si layer and the inset image demonstrates the photon absorption profile in the 1.2 μm height device [16]

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that the optical absorption and Jsc are increased monotonically with NSP height. The inset image in Fig. 2f illustrates the absorption profile of the device based on NSP with 1.2 μm height, indicating that the most of absorption occurred in the active a-Si active layer. Moreover, it can be concluded that by increasing the pitch size for the same height, the optical absorption is decreased due to lower light scattering with large NSP to NSP distance in large pitch [16].

3  Anti-reflection Nanostructures for Solar Cell Devices Enhancing the  Light absorption  of Anti-reflection (AR) nanostructures the active layer  in the optoelectronic devices is necessary step to boost the device performance. To satisfy this purpose, there are two general strategies, first, increase the optical absorption and decrease the reflectance by using the anti-reflection (AR) layer and second, increase the absorption through the  surface  of nanostructure. Among AR methods, anti-reflection coating is a common way, which is typically utilized for enhancement of light absorption. However, this technique needs several coating layers to achieve a broad-band absorption, results in higher production cost. For the absorption enhancement methods, random nanotextured surface and back metal reflector are usually employed. In fact, the absorption can increase up to the Lambertian limit of 4N2, where N is the refractive index [20, 21]. Basically, when light passes through a textured media with a roughness of z0 (peak to valley), there is a phase shift proportional to n1 · z + n2 · (z0−z). Here, z is corresponded to the distance from the highest point to the traveled interface in first medium, which has a refractive index of n1 · (z0−z) is the traveled distance after interface in medium 2 with refractive index of n2, as can be found in Fig. 3. Due to having textured interface, the z is depended on the surface morphology and geometry, following the lateral coordinates of x and y. As a result, the phase shift of a plane wave upon passing a rough interface is defined by (n1−n2) · z(x, y) + n2 · z0, indicating the effect of textured surface on the plane wave [22].

Fig. 3  Acquired phase shift by a plane wave after passing a textured interface [22]

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In fact, for an ideal Lambertian scatterer with an isotropic radiance, the value of radiance (L) is equal to 1/π. However for an arbitrary scatterer, the value of L is depended on the refractive indices of both medias and the surface morphology (z(x, y)) of the interface. For example, in a silicon solar cell, consisted of a glass substrate with refractive index of ~1.5, a transparent electrode (n ~ 2) and the silicon absorber (n ~ 4), by texturing the electrode/silicon interface, light scattering can be achieved. Here, the scattering of light from the electrode to air is so close to Lambertian and more importantly the fraction of the scattered light shows an angular dependence with respect to the Lambertian scatterer. In order to consider scattering into the silicon layer, it is proved that the height (Z) has to be scaled by1/2 and the lateral dimensions should be scaled by 1/4. Consequently, for the case of Lambertian scatterer to silicon, it is required to do a reduction in feature size by a factor of 4 and an enhancement in aspect ratio by a factor of 2. In general, when the light passes from an arbitrary interface with refractive indices of n1 and n2 to another interface with refractive indices of n1′ and n2′ , we can consider the following scaled coordinates [22–24]:

x′ = x ·n2 / n2′ ; y ′′ == y· n2 / n2′ ; z ′ = z·| n1 − n2 | / | n1′ − n2′ |

In fact, the scaling laws is really important, and it can provide an intuitive estimation of the desired scattering profile by passing the light from different interfaces. In order to further improve the light scattering close to ideal Lambertian scatterer, the aspect ratio of the textured media needs to be increased [25]. However, this approach is detrimental for the fabrication process and electrical properties of the devices and thus there is a trade-off here [26, 27]. If we can consider the 4n2 Lambertian limit as 2 × 2 × n2, the first 2 is ascribed to the back-reflector effect in a solar cell, which increases the light path twice. The second 2 can be corresponded to the Lambertionality factor (the average light path enhancement). The n2 factor is proportional to 1/b, where b is the fraction of light, which escapes from the textured interface [22]. In order to further enhance or exceed this limit, periodic three-dimensional (3-D) nanostructures have been employed. It was discovered that the effectiveness of light absorption by using these nanostructures is related to the materials as well as geometry. In case of geometry, when the nanostructure is larger than the optical wavelength, enhanced optical travel path and higher absorption are achieved due to better light scattering inside the nanostructure, whereas, the mechanism of light absorption for nanostructure with subwavelength geometry can be fundamental photonic resonant modes due to light confinement. In a silicon solar cell, nanotextured surface can reduce the light reflection; however, it can increase the Auger recombination as well. Since nanotextured surface increases the surface area significantly, proper passivation approaches are required to decrease the recombination. To address recombination issue of nanocone in silicon solar cells, some people design an emitter at the back of the device rather than its top side [21]. As shown in Fig. 1a, b, the back of the device has highly p+ and n− doped regions, whereas the front side is consisted of the nanocone array, which

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turns the device color to black. Using this structure, impressive EQE values of >80% was achieved over the spectrum from 400 nm to 800 nm, as shown in Fig. 4c, due to the presence of nanocone array [21]. In fact, the maximum light absorption is achieved by nanocone at the front side and minimum PV losses is achieved at the back side of the device, resulting in 30.7% enhancement of JSC for textured device. The photovoltaic curve of the corresponding device showed significant improvement of the Jsc for the textured device as compared to the planar one without losing any Voc. The PCE for the nanocone-based device was 13.7%, which was much higher than the planar one with 10.9%, indicating the advantage of the nanostructures [21]. This mechanism is dominant for nanostructures such as nanocone and nanospike with a gradient of the effective refractive index, which can work as broad-band and omnidirectional AR layer [21, 28]. This mechanism is also observed for uniform diameter of nanostructures such as nanowell, nanowire, and nanopillar. In these cases, the diameter and length play a significant role in optical absorption due to plasmonic resonance which is another effective light management scheme [29]. Plasmonic effect is normally applied in a solar cell device using metallic nanoparticles. These nanoparticles showed surface plasmons, which is corresponded to the excitations of the electrons at a metal/dielectric interface [29, 30]. By proper designing of the morphology and geometry of these metallodielectric structures, light can be focused on absorber layer of the device, resulting in higher absorption. In fact, both plasmonic effects from the localized surface plasmon in the metallic nanoparticles and the surface plasmon polaritons (SPPs) at the interface of metal/semiconductor are beneficial for increasing the absorption [29–32]. Plasmonic effect can reduce the thickness of the absorber layer while maintaining the same optical absorption in three ways, as shown in Fig.  4e. In the first approach, the metallic nanoparticles are employed as scattering elements for subwavelengths to trap the plane waves from the sunlight into the absorber layer. In the second way, they can be employed as subwavelength antennas, where the plasmonic field is coupled with the absorber layer to increase the absorption. In the third approach, a corrugated layer of metallic film is formed at the back of the device, which can couple the plane waves of lights into the SPP modes at the metal/absorber interface and conduct the modes into the absorber layer [29]. Then, the scattered light obtains an angular spread in the absorber layer and increases the optical path length, resulting in light trapping. In fact, the light beams will pass many times in the absorber layer, enhancing the effective path length [30, 31]. Additionally, the size and shape of metallic nanoparticles also play critical roles for improving the incoupling efficiency [32]. Figure 4f shows that the smaller size of metallic nanoparticles (close to the semiconductor layer) can couple larger amount of light into the absorber layer due to the improved near-field coupling.

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Fig. 4 (a) Photographs of the back and front of the 10-mm-thick Si solar cell. Inset images are the SEM images of nanocones with different magnifications. Scale bars are 2 mm, 5 mm, and 400 nm for top, bottom-left, and bottom-right, respectively. (b) Schematic of the silicon solar cell. (c) EQE spectra, and (d) J–V curves of the silicon solar cells without and with nanocone [21]. (e) Scattering of the light by metal nanoparticles on the surface of device, resulted in light trapping (left), trapping of the light caused by the excitation of localized surface plasmon in metal nanoparticles located inside of the semiconductor (middle), and trapping of the light by the excitation of surface plasmon polaritons located at the interface of meta//semiconductor. (f) Portion of scattered light inside substrate with respect to the total scattered power for the silver nanoparticles on silicon with different sizes and shapes. (g) Enhanced electrical field intensity by metal nanoparticles close to the surface [29]

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Figure 4g shows the light concentration around the surface of metallic nanoparticles due to resonant plasmon excitation, which significantly increases the local field around the nanoparticle and thereby the light absorption. Therefore, the metallic nanoparticles work as an effective antenna for the sunlight and enhance the localized surface plasmon mode [29, 32]. In another example, the role of nanocone arrays as an anti-reflection layer is studied in a CdS/CdTe solar cell. Figure 5a shows the angular view of the nanocone array made of PDMS. In order to fabricate these nanostructures, the electropolished aluminum foil is imprinted by a silicon mold containing nanopillars. Then, several steps of anodization and etching are necessary to obtain an ordered inverted nanocone structure of AAO, as reported elsewhere [33]. Afterward, the AAO template is used as a mold and the PDMS solution is casted on this template, followed by heating and peeling of the nanocone PDMS AR layer. After preparation of PDMS AR layer, it can be attached to the glass side of device without any glue, as shown in Fig. 5b. The PDMS film has a strong Van der Waals interaction with glass substrate and thus it can be easily attached or detached from the devices. This self-­assembly property leads to convenient mounting and replacement of these AR films which is promising for industrial applications [34]. Interestingly, these nanostructures not only increase the optical absorption but also have superhydrophobicity properties as illustrated in Fig. 5c. As it can be observed, nanocone can increase the contact angle of water droplet from 98° to 152°. This characteristic of the nanocone array can increase the stability of a solar cell device in a humid environment, which is beneficial for practical application [33]. This nanocone AR film can be attached to different types of solar cell devices. For instance, the simulated reflectance spectra of solar cells without and with AR film from 400 to 900 nm are shown in Fig. 5d. In this measurement, both pitch and height of the nanocone were 1 μm. By applying AR film, the reflectance of film is decreased by 4% as compared to device without AR film, which is due to the tapered nanocone geometry inducing a gradual change in refractive index from air to PDMS [33]. Moreover, the difference between refractive indexes of glass and PDMS is negligible. After applying AR film, there is still a bit reflectance which comes from the interface of FTO layer and electron transport layer (ETL). This can be addressed by roughening the surface of glass before FTO coating. The cross-sectional electric field intensity (E2) distribution of electromagnetic wave at 600  nm calculated by FDTD simulation is shown in the inset images of Fig. 5d. Notable, the electromagnetic plane wave propagates downward. As seen, in the presence of nanocone, the interference and reflected wave are weaker and the electric field distribution inside the nanocone structure is stronger [35]. As mentioned earlier, the geometry of nanocone (pitch and height) can be largely tuned based on the fabrication method. The optimization of nanocone geometry is shown in Fig. 5e. For this purpose, the total integrated reflectance was estimated and plotted versus pitch and height of nanocone structure. The results show that PDMS nanocone with 1 μm pitch and 1 μm height is the optimum structure to reduce the reflectance significantly. As a result, the external quantum efficiency (EQE) spectra of CdS/CdTe solar cells without and with AR film, shown in Fig. 5f, clearly demonstrates lower reflectance for device

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Fig. 5  Angular-view SEM image (inset is top-view) of (a) inverted nanocone made of nanocone array of PDMS AR film. (b) Schematic of CdS/CdTe solar cell device with AR film on top. (c) Contact angle measurement of nanocone array and planar PDMS films using a droplet of water. (d) FDTD simulations of reflectance spectra for the corresponding devices. The inset images show the cross-sectional distribution of electric field intensity for electromagnetic wave at 600 nm wavelength. (e) Optimization of nanocone geometry using reflectance calculation by FDTD simulation. (f) External quantum efficiency measurement of CdS/CdTe solar cells without and with PDMS nanocone AR layer. The inset graph is the reflectance measurement obtained from different light incident angles from 0° to 60° for CdS/CdTe solar cell without and with AR layer [33]

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with PDMS nanocone AR film. For CdS/CdTe device, the Jsc improvement is almost 1.11  mA/cm2 (4.6% enhancement) [33]. It is worth pointing out that the light absorption in device with nanocone AR layer for angular incidence is significantly improved as compared to device without AR film. The inset graph in Fig. 5f depicts the reflectance spectra of CdS/CdTe solar cells without and with PDMS nanocone for light incident angles changed from 0° to 60° using an integrating sphere. From these results, the nanocone sample remains invariant. The measured Jsc under this condition also confirmed the advantage of nanocone AR film. In nanocone sample, the Jsc enhancement by changing the angle from 0° to 60° is increased as compared to planar one indicating the advantage of nanocone AR film particularly at larger angles. This characteristic from nanocone array is beneficial for the solar cells installed in a solar farm due to the variation of incident sun light during the daytime and different seasons [33]. The AR films can be used to improve the optical absorption in different types of solar cell devices. Organo-metal halide perovskite materials have attracted tremendous attention of many research groups due to their interesting properties such as low-cost fabrication and materials, long carrier diffusion length, high mobility, and high optical absorption [36]. These properties make them ideal candidates for fabrication of highly efficient solar cell devices with PCE of over 25% [37]. There are a number of different processes reported for the fabrication of perovskite films such as spin-coating, deep-coating, vacuum deposition, spray coating, and chemical vapor deposition [38]. These fabrication methods require low temperature and thus this material has a great potential for fabrication of flexible solar cells. The application of AR film on perovskite solar cell has been also studied to decrease the reflectance and further improve the device performance. The absorption and reflectance spectra of perovskite films (400 nm-thick) without and with nanocone AR films based on different aspect ratios can be found in Fig. 6a. In fact, by applying the AR film, the reflectance is decreased and as shown, this reduction is the highest for nanocone with aspect ratio of 1, indicating the importance of geometry optimization. The integrated absorption of perovskite films without and with AR film versus different angles of the incident light indicates unchanged light absorption for nanocone-­ based device at high angles as compared to device without nanocone (Fig. 6b). In addition, the contact angle measurement of water droplet on top of nanocone array demonstrates the excellent superhydrophobicity property of nanocone AR film with aspect ratio of 1.0 (Fig. 6c). This indicates that by inserting nanocone AR film, the perovskite solar cells can have self-cleaning property, which is beneficial for industrial applications (Fig. 6d). Since the perovskite films are highly sensitive to moisture, these AR films can increase the stability of devices, which is the main challenging issue for this type of solar cell [39]. The FDTD simulations for perovskite solar cells without and with nanocone AR film are shown in Fig. 6e, f. The generation rate profile in both cases can be observed in the cross-sectional views of the perovskite films, where the electromagnetic plane waves propagated upward to simulate the real situation. The focal point in the film with nanocone AR film (reddish colored zone) represents the region with a high

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Fig. 6 (a) Absorption and reflectance spectra of perovskite film before and after applying nanocone AR layers with different aspect ratios. (b) Angular absorption measurement of perovskite film without and with nanocone array with aspect ratio of 1. (c) Contact angle measurement of water droplet on top of PDMS nanocone array with different aspect ratios. The inset image is the micrograph of water droplet on AR film with aspect ratio of 1. (d) Self-cleaning experiment on perovskite solar cells with (d1, d2) and without (d3, d4) applying AR film. FDTD simulation of perovskite solar cells without (e) and with (f) AR film indicating the generation rate (number of absorber photons/ (m3.s)) inside the perovskite absorber layer [39]

generation rate of electron–hole pairs. Notably, hot zone with high generation rate is not obvious in device without AR film [40]. Figure 7a shows cross-sectional view SEM image of perovskite solar cell fabricated on flexible glass using a two-step evaporation technique. As shown, the device consisted of ITO as an electrode, ZnO as an electron transporting layer (ETL), perovskite absorber, and spiro-OMeTAD as a hole transport layer (HTL), and gold as a back contact. The effect of nanocone AR film on J-V characteristic of perovskite

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Fig. 7 (a) Cross-sectional view SEM image of perovskite solar cell fabricated by a two-step evaporation method, which is made of ITO coated willow glass, ZnO ETL, perovskite absorber, Spiro as HTL and gold contact. The inset image is SEM image of PDMS nanocone array as AR layer, which is stacked on the willow glass. (b) J-V curve of perovskite solar cell without and with AR film. The inset is schematic of device with AR film. (c) Jsc and PCE measurements of device without and with AR film under shining incident light with different angles [39]

solar cell has been explored in Fig. 7b. From the result, it is obvious that by inserting the AR film, the current density of device is increased up to 1.6  mA/cm2. More interestingly, when the light incidence is increased, both Jsc and PCE of device without AR film drop faster than those of for device with AR layer. Particularly, the PCE of device with AR film is enhanced by 20% at 60° incident angle as compared to the reference sample (Fig.  7c). It is worth pointing out that the nanocone array has omnidirectional light harvesting capability and this property is beneficial for solar panels without mechanical solar tracking system [39].

4  Nanotextured Substrates for Solar Cell Devices The clean and renewable solar energy harvesting have gained considerable attention due to the presence of new and promising technologies, which are highly attractive for low-cost applications and low temperature processing as compared to crystalline silicon devices. Efficient light absorption in thin film technology is significant for  the fabrication of highly efficient solar cells. As mentioned in the previous

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sections, 3D nanostructures have been extensively studied to boost the performance of thin film devices utilizing various advanced light management schemes [40]. As discussed in the previous section, these nanostructures can serve as an anti-­ reflection film to reduce the light reflection, resulting in higher performance. Additionally, they have superhydrophobicity and self-cleaning properties, which can increase the stability of devices. In addition to this application, these nanostructures can be used as substrates for the fabrication of photovoltaic devices. For this purpose, all layers in the device need to be deposited conformally on the surface of nanostructures with the minimum level of the defects. Otherwise, there is no advantage using nanostructures as substrates for solar cell device with many defects. There are a variety of 3D nanostructures such as nanocones, nanowires, nanopillars, nanorods, nanopyramids, nanospikes, nanospheres, and nanowells [41]. Among these nanostructures, nanocones show promising capabilities such as enhanced sunlight harvesting in different incident angles and improved mechanical properties of devices which are applicable in flexible solar cell devices [42]. Figure 8a1–a4 demonstrate the fabrication process of a-Si solar cell on polyimide (PI) substrate, where the inverted nanocone AAO template was considered as a mold and the PI substrate was obtained after casting of the PI solution on top of the AAO mold. For the fabrication of a-Si solar cell on top of PI nanocone, all layers were deposited using vacuum techniques. The angular-view SEM images of textured devices with aspect ratios of 0.5 and 1.0 are shown in Fig. 8b, c, indicating a very uniform deposition of all of the layers [43]. The J-V curves of the corresponding devices indicate that the nanocone substrates increase the current density and fill factor of a-Si solar cells due to higher absorption and better charge collection (Fig. 8d). Moreover, the aspect ratio of 0.25 demonstrates the best results, suggesting the most uniform layers with lowest defect levels. As illustrated in Fig. 8e, the EQE of the corresponding devices are in good agreement with J-V results [43]. Notably the nanocone structures sacrificed the open circuit voltage (Voc) of the devices as compared to flat substrates due to having more interface recombination. This suggests that the nanocone samples are induced surface recombination sites due to the larger surface area at the interface. The Voc drop is the highest value for nanocone sample with aspect ratio of 1.0 due to the difficulty in  the fabrication process in order to achieve a very conformal films. As mentioned in the last section, nanostructures AR film can improve the optical absorption of devices in high incident angle such as 60°. This advantage is also demonstrated in the device fabricated on nanostructures. For instance, nanocone-based device shows much higher integrated absorption and thus PCE as compared to the flat substrate with relatively 15% enhancement (Fig. 8f, g) [43]. Also, the nanocone device depicts much better performance in all incident angles as compared to the flat device, due to the gradual change in the effective refractive index of the entire nanostructures. This omnidirectional enhancement of the absorption and performance for the nanocone device is of significance for industrial applications without applying costly solar tracking system. In addition to the above discussion, nanostructured devices have better mechanical properties as compared to flat devices, which make them ideal candidates for flexible devices (Fig. 8h, i)

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Fig. 8 (a1–a4) Fabrication process of a-Si solar cell on plastic nanocone substrate. Cross-section and angular SEM images of a-Si solar cells fabricated on nanocone plastic substrates with aspect ratios of (b) 0.5 and (c) 1.0. (d) J-V curves and (e) EQE spectra of a-Si solar cells fabricated on planar and nanocone arrays with aspect ratios of 0.5 and 1.0. (f) Broad-band integrated absorption of a-Si absorber and (g) PCE of devices deposited on planar and nanocone substrates by tuning the incident angles. Efficiency of a-Si solar cells on flat and nanocone array upon increasing the bending angle (h), and bending cycles (i) [43]

[43]. As seen, the PCE of nanocone devices remains almost the same by tuning the bending angle due to lower surface reflection and better mechanical property of nanocone sample as compared to flat device. In addition, by performing 1000 bending cycles, the nanocone device retained 95% of its initial PCE value, which is much better than the flat device with 36% PCE loss. This clearly showed the beneficial advantages of nanocone arrays as substrates to improve the mechanical properties of flexible a-Si solar cell [43]. The application of nanostructure substrates has been also demonstrated for perovskite solar cell devices. Nanotube structure of TiO2 could be an ideal electron transporting layer for the perovskite solar cell. One of the great approaches to

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synthesize TiO2 nanotube is anodization of the titanium foil. Using this technique, not only Ti is converted to the oxide but also the oxide layer will be in the form of nanowells structure. Another advantage of this nanostructure is the flexibility of the substrate [44]. Over the past years, several groups focused on this direction and fabricated high quality TiO2 nanotube and demonstrated their application in the perovskite solar cells. Figure 9a shows the schematic of the corresponding device with carbon nanotube as top electrode. Cross-sectional SEM image of this device is shown in Fig. 9b, indicating perovskite pillar inside the nanowells with ~300 nm length. In order to improve the photovoltaic properties of the device, they optimized the length of TiO2 nanotube and treated them with aqueous solution of TiCl4 [44]. Figure 9c depicts the J-V curves of the corresponding devices, indicating PCE of 8.31%. They found that 25 micron-length of nanotube with TiCl4 provides the highest efficiency with much lower voltage loss. Moreover, the calculated Jsc from the EQE spectra and solar spectrum for the corresponding devices are well-matched with the J-V results. They also demonstrated the excellent flexibility of the device as can be seen in the inset photograph of Fig. 9d [44].

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Another example is the application of inverted nanocone substrates for the fabrication of photovoltaic devices with different aspect ratios. In this regard, inverted nanocone AAO was employed as a mold and the PDMS nanocone film was peeled off from the template after casting to obtain another mold. Then, the epoxy solution was casted on top of the PDMS nanocone, followed by UV curing. The obtained inverted nanocone epoxy was used as a substrate for solar cell fabrication (Fig. 10a1– a2). The top and angular view SEM images of PDMS nanocone mold and inverted nanocone epoxy substrate are depicted in Fig. 10b, c, indicating very uniform and ordered nanostructures. Figure  10d1–d4 demonstrate the fabrication process of perovskite solar cell using a two-step evaporation technique for perovskite deposition. The generation rate of electron–hole pairs in the perovskite devices fabricated on inverted nanocone substrates with different aspect ratios was studied using FDTD simulation as illustrated in Fig. 10c1–c4. As seen, the aspect ratio of 1.0 indicates the highest generation rate specially in the proximity of the inverted nanocone tip. This can be related to the lowest reflection in the sample with high aspect ratio as compared to the planar device (Fig. 10f). The J-V results for the corresponding devices indicate that the aspect ratio of 0.5 is the best substrate for device fabrication in terms of both uniformity and optical absorption. As shown in Fig. 10g, the Jsc and PCE enhancements for device with aspect ratio of 0.5 are 37% and 38% as compared to the planar device, respectively [45]. These improvements are lower for the aspect ratio of 1.0 due to ununiform layers deposited on inverted nanocone structures. It is noteworthy to mention that the textured devices demonstrate lower Voc as compared to the planar one. This could be related to larger surface area in nanostructure samples, inducing more recombination sites [45]. As discussed earlier, the nanostructures can improve the mechanical properties of solar cell devices, which is beneficial for flexible photovoltaic devices. The mechanical properties of perovskite solar cells fabricated on inverted nanocone were tested using a bending setup (Fig. 11a). The textured device after 200 bending cycles maintained 95% of its initial PCE value which is higher than the  planar device with 30% PCE loss, as observed in Fig. 11b. The top-view SEM images of the corresponding devices suggest that after 200 bending cycles, many crack lines existed in the planar samples due to the difference between Young’s modulus (E) for different layers of the device. However, the inverted nanocone device showed no crack line indicating the role of nanostructure in improvement of the mechanical properties for the nanotextured devices (Fig. 11c, d). Moreover, the finite element modeling of the corresponding device can help to have a better understanding of the cross-sectional stress distribution through the devices. Figure 11e, f depict the contour images of stress distribution for planar and textured devices upon applying 10 N force at the edge of samples, calculated by COMSOL Multiphysics software. As seen, the concentrated stresses for planar device existed between different layers of the device. Interestingly, the nanotextured device annihilates these concentrated stresses at the interfaces and makes a uniform distribution of stress within all the layers, resulting in better mechanical properties and flexibility for the devices [43, 45].

Fig. 10 (a1–a4) Fabrication process of inverted nanocone AAO, nanocone PDMS, and inverted nanocone epoxy. Top and angular view SEM images of (b) nanocone PDMS mold and (c) inverted nanocone epoxy substrate. (d1–d4) Fabrication process of perovskite solar cells on the inver­ ted nanocone epoxy. (e1–e4) FDTD simulation (Generation rate) of perovskite solar cells fabricated on planar and inverted nanocone substrates with aspect ratios of 0.25, 0.5, and 1.0. (f) Reflectance spectra of perovskite films and (g) J-V curves of perovskite solar cells fabricated on inverted nanocone epoxy with aspect ratios of 0 (planar), 0.25, 0.5, and 1.0. The inset image is cross-sectional view SEM image of perovskite solar cell on inverted nanocone with aspect ratio of 1.0 [45]

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Fig. 11 (a) Bending cycle measurement setup, (b) PCE and Jsc measurements of perovskite solar cells fabricated on planar and inverted nanocone epoxy upon doing bending cycles. Top-view SEM images of perovskite films on planar (c) and inverted nanocone (d) epoxy substrates. Finite element mechanical modeling of planar (e) and textured (f) substrates, indicating the distribution of stress inside the device structure [45]

5  Conclusions The use of nanostructures for light management purpose has a great potency to boost the solar to electric power conversion efficiency in the solar cell devices. Using a nanostructure, the absorber layer can harvest the sun light with thinner thickness as compared to planar device due to the improvement of carrier collection

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and light tapping property. Nanostructures can be served as an anti-reflection layer or a substrate for solar cell devices. In case of anti-reflection film, plastic nanostructures have great superhydrophobicity and self-cleaning properties, which makes them excellent candidates for the fabrication of water-repellent and stable solar cell devices. Moreover, these nanostructures improve the optical absorption of device at large light incident angles. These properties for nanostructured-based devices are beneficial for industrial applications in a solar farm. For this application, nanocone arrays have the best geometry for light absorption and device fabrication purposes as compared to other nanostructured counterparts. For nanotextured substrates used in photovoltaic devices, there are more advantages as compared to the planar device such as enhanced carrier collection and improved mechanical properties. In fact, nanostructures strongly fold all of the layers in the device and greatly annihilate the concentrated stresses particularly at the interfacial regions of the device. This architecture design can enhance the robustness and mechanical flexibility of thin film solar cells and applicable for most of solar cells.

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Nanomaterials Enhanced Heat Storage in Molten Salts Xiaotong Guo, Di Hu, Linpo Yu, Lan Xia, and George Z. Chen

Abstract  Molten salts have great potentials as thermal energy storage (TES) media due to their many types and compositions, wide operating temperatures and relatively high values of specific heat capacity. The rapid development of TES technologies in recent years has triggered the modification of TES media to higher operating temperatures and hence larger heat capacities for capturing and storing heat from a wider range of sources, especially renewables and thermal wastes from industry. In doing so, many nanomaterials have been added to molten salts, particularly their eutectic mixtures, as heat storage enhancers as reported in the literature, achieving remarkable increases in specific heat capacity. To explain the mechanism of the resultant capacity variations, different factors of the nanomaterials including particle size, concentration and structure, as well as interaction between the nanomaterials and the molten salts have been investigated with divergent results. It has been speculated that the formation of nanolayers on the surface of the nanomaterials and the induced larger surface area may be responsible for the enhanced specific heat capacity of molten salts. Besides, the principle of formation of the nanostructures and the specific heat capacity prediction model were also proposed. More efforts are needed to systematically analyse and integrate previous findings, including those contradictions originated from the scattered methods and techniques applied, to facilitate more comparable and repeatable outputs in the further work. Keywords  Thermal energy storage · Specific heat capacity · Molten salts Nanomaterials

X. Guo International Doctoral Innovation Centre, Faculty of Science and Engineering, University of Nottingham Ningbo China, Ningbo, People’s Republic of China D. Hu · L. Yu · L. Xia Department of Chemical and Environmental Engineering, Faculty of Science and Engineering, University of Nottingham Ningbo China, Ningbo, People’s Republic of China Energy Engineering Research Group, Faculty of Science and Engineering, University of Nottingham Ningbo China, Ningbo, People’s Republic of China © National Technology & Engineering Solutions of Sandia, LLC 2021 M. Alston, T. N. Lambert (eds.), Energy-Sustainable Advanced Materials, https://doi.org/10.1007/978-3-030-57492-5_6

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1  Introduction An increasing percentage of renewable energy is being deployed in the total amount of global energy supply in light of the challenging criteria and scenarios for the reduction of emissions of greenhouse gases from using fossil fuels. However, the biggest drawback of renewable energy is the intermittence of harvest and conversion, most of which is quite time-and-climate dependent. To mitigate or work on  these issues, various materials and devices have been designed to store the excess energy in different forms, such as mechanical storage in flywheels, electrochemical storage in batteries, and thermal energy storage (TES) in phase change materials (PCMs), solids or molten salts. (Note: Heat is defined as thermal energy in transit, which is a process function in a thermodynamic system. In this chapter, thermal energy and heat are used as exchangeable terms to which no depiction of a particular process is referred). Figure 1 shows a brief working process of heat source utilisation with the integration of TES [1–3]. Commonly, the TES systems are classified by storage concept (active or passive) and storage mechanism (sensible, latent or chemical/thermochemical) [4]. In the sensible heat storage, sensible TES materials (e.g., steam, oils and molten salts) undergo no changes in phase over the operating temperature range [1]. Latent heat storage using phase change materials (PCMs) works in a nearly isothermal way during melting/solidification or gasification/liquefaction processes whilst storing or releasing thermal energy [5]. Thermochemical energy storage utilises sorption materials. For example, silica gel, magnesium sulphate, lithium bromide, lithium chloride and sodium hydroxide are all capable of absorption or adsorption of water vapour. Since the sorption process of reactive components is thermally reversible, high heat storage capacity can be achieved with negligible thermal losses during the storage period [3]. In addition, the TES system can also be categorised into low, medium and high temperature TES systems by their operating temperatures. Low-temperature TES takes advantages of lower rates of off-peak electricity based on a short-term storage capacity in a day. The storage medium is working at operating temperatures comparable with the spatial temperature in heating/cooling applications. Chilled water, phase change materials (PCMs), ice or cryogen (e.g., liquid air and liquid nitrogen) are common options. Medium-temperature TES usually operates at the temperature slightly higher than the spatial temperature range in the facilities such as solar hot water, air heating and solar heating/cooling units in buildings and G. Z. Chen (*) Department of Chemical and Environmental Engineering, Faculty of Science and Engineering, University of Nottingham Ningbo China, Ningbo, People’s Republic of China Energy Engineering Research Group, Faculty of Science and Engineering, University of Nottingham Ningbo China, Ningbo, People’s Republic of China Department of Chemical and Environmental Engineering, Faculty of Engineering, University of Nottingham, Nottingham, UK Advanced Materials Research Group, Faculty of Engineering, University of Nottingham, Nottingham, UK e-mail: [email protected]

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Solar DNI Geothermal

Heat Source

Power Block

Direct full/partial load supply

Fuels Electricity etc.

Waste heat etc. Charging process

Discharging process

Thermochemical

Latent heat

Sensible heat

Passive storage

Active storage

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Fig. 1  A process sketch of heat source utilisation with the integration of TES (DNI: Direct Normal Irradiance. DSG: Direct Steam Generation) [1–3]

industry [1]. High-temperature TES is especially suitable to be integrated into concentrating solar power (CSP) plants to solve the mismatch between energy supply and demand. Since 2006, operators have built TES systems into CSP plants almost exclusively using sensible heat for heat storage [2]. In the TES system, the design of the storage configuration depends on the scaled capacity and storage media. Therefore, for a given scale of storage capacity, TES media with higher heat capacity can reduce the amount of materials required, resulting in smaller size of storage tanks and less pumping power requirement, and consequently reduce the overall installation and maintenance costs. Molten salts referring to salts which are heated to the liquid phase have been regarded as promising TES materials. They have advantages such as low vapour pressure, high operating temperature, and good thermal stability, which makes them appealing for high-temperature sensible heat storage. However, the specific heat capacity of a pure molten salt or salt mixture is often