Nanostructured Materials: Applications, Synthesis and In-Situ Characterization 9783110458299

Table of contents :
Cover
Half Title
Also of interest
Nanostructured Materials: Applications, Synthesis and In-Situ Characterization
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Contents
List of abbreviations
List of contributors
1. Introduction
References
2. Guest-induced luminescence of metal-organic frameworks
2.1 Introduction
2.2 Synthesis and loading methods
2.2.1 MOF synthesis
2.2.2 Encapsulation of guest species
2.3 Materials used for guest-induced MOF luminescence
2.3.1 MOF encapsulation of lanthanide ions
2.3.2 MOF encapsulation of metal ion complexes
2.3.3 MOF encapsulation of quantum dots
2.3.4 MOF encapsulation of plasmonic nanoparticles
2.3.5 MOF encapsulation of organic phosphors
2.4 Applications of guest-induced luminescence in MOFs
2.4.1 White light emission
2.4.2 Chemical sensing of ions
2.4.3 Other applications
2.5 Concluding remarks
References
3. Synthesis, functionalisation, and applications of iron oxide magnetic
3.1 Introduction
3.2 Types of iron oxide magnetic nanoparticles
3.3 Synthesis methods
3.4 Advanced studies on iron oxide formation
3.5 Applications
3.6 Surface functionalisation
3.6.1 Coating with luminescent materials
3.6.2 Coating with metal–organic frameworks (MOFs)
3.7 Conclusions and outlook
References
4. Nanostructured bioceramics
4.1 Introduction
4.2 Types of nanostructured bioceramics
4.2.1 Unique properties of medically relevant nanobioceramics
4.2.2 Calcium-phosphate-based bioceramics
4.2.3 Calcium-carbonate-based bioceramics
4.2.4 Glass-based bioceramics
4.2.5 Porous materials
4.3 Synthesis methods for nanostructured bioceramics
Co-precipitation
Sol-gel processing
4.4 Functionalisation of bioceramics
4.5 Applications of bioceramics
4.6 Conclusions and outlook
References
5. In situ luminescence analysis of coordination sensors (ILACS): looking inside chemical reactions
5.1 Introduction
5.2 Working principle
5.3 Spectroscopic background of coordination sensors
5.4 Examples of multimodal in situ experimental instrumentation
5.5 Selected application examples
5.6 Tracking the reaction pathways without coordination sensors
5.7 Conclusions and outlook
References
6. In situ monitoring of the syntheses, phase transformations, and loading processes of metal–organic frameworks
6.1 Introduction
6.2 In situ analysis of MOF synthesis and phase transformations
6.2.1 In situ XRD analysis of solution-based processes
In situ XRD analysis employing batch reactors
In situ XRD analysis employing flow reactors
6.2.2 In situ XRD analysis of solvent-free processes
6.2.3 In situ IR spectroscopy
6.2.4 Other analysis techniques
In situ UV/Vis absorption spectroscopy
In situ PDF measurements
In situ light scattering
In situ atomic force microscopy (AFM) studies
6.3 In situ structural analysis of MOF loading processes
6.3.1 Gas storage
Hydrogen
Carbon dioxide
Oxygen
6.3.2 Removal of harmful molecules
6.3.3 Breathing and gate-opening effects
6.3.4 Negative gas adsorption (NGA) transitions
6.4 Loading processes on MOF thin films
6.5 Concluding remarks
References
7. Operando studies on the charge and discharge processes of battery materials
7.1 Introduction
7.2 Theoretical background: intercalation and conversion reaction batteries
7.2.1 Intercalation reaction mechanisms
7.2.2 Conversion reaction mechanism
7.3 Operando X-ray diffraction
7.3.1 Lithium-ion batteries (LIBs)
7.3.2 Sodium-ion batteries (SIBs)
7.3.3 Potassium-ion batteries (PIBs)
7.3.4 (Hybrid) zinc-ion batteries (ZIBs)
7.3.5 Magnesium-ion batteries (MIBs)
7.3.6 All-solid-state batteries (ASSBs)
7.4 Operando X-ray absorption spectroscopy
7.5 Operando imaging techniques
7.6 Operando optical and vibrational spectroscopy
7.7 Other operando techniques
7.8 Conclusions and Outlook
References
Index

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Huayna Terraschke (Ed.) Nanostructured Materials

Also of interest Chemistry of Nanomaterials. Vol.  Multifunctional Materials S.S.R. Kumar Challa (Ed.),  ISBN ----, e-ISBN (PDF) ----, e-ISBN (EPUB) ---- Chemistry of Nanomaterials. Vol.  A Metallic Nanomaterials (Part A) S.S.R. Kumar Challa (Ed.),  ISBN ----, e-ISBN (PDF) ----, e-ISBN (EPUB) ---- Chemistry of Nanomaterials. Vol.  B Metallic Nanomaterials (Part B) S.S.R. Kumar Challa (Ed.),  ISBN ----, e-ISBN (PDF) ----, e-ISBN (EPUB) ---- Nanoparticles Raz Jelinek,  ISBN ----, e-ISBN (PDF) ----, e-ISBN (EPUB) ----

Structure Analysis of Advanced Nanomaterials. Nanoworld by High-Resolution Electron Microscopy Takeo Oku,  ISBN ----, e-ISBN (PDF) ----, e-ISBN (EPUB) ----

Nanostructured Materials Applications, Synthesis and In-Situ Characterization Edited by Huayna Terraschke

Editor Jun.-Prof. Dr. Huayna Terraschke Institut für Anorganische Chemie Christian-Albrechts-Universität zu Kiel Max-Eyth-Str. 2 24118 Kiel Germany [email protected]

ISBN 978-3-11-045829-9 e-ISBN (PDF) 978-3-11-045909-8 e-ISBN (EPUB) 978-3-11-045833-6 Library of Congress Control Number: 2023937882 Bibliographic information published by the Deutsche Nationalbibliothek The Deutsche Nationalbibliothek lists this publication in the Deutsche Nationalbibliografie; detailed bibliographic data are available on the internet at http://dnb.dnb.de. © 2024 Walter de Gruyter GmbH, Berlin/Boston Cover image: Huayna Terraschke Typesetting: Integra Software Services Pvt. Ltd. Printing and binding: CPI books GmbH, Leck www.degruyter.com

Contents List of abbreviations List of contributors

VII XIII

Katia Nchimi, Huayna Terraschke Chapter 1 Introduction 1 Moritz Bassen, Katia Nchimi, Huayna Terraschke Chapter 2 Guest-induced luminescence of metal-organic frameworks

5

Patric Lindenberg, Katia Nchimi, Huayna Terraschke Chapter 3 Synthesis, functionalisation, and applications of iron oxide magnetic nanoparticles 47 Niklas Ruser, Huayna Terraschke Chapter 4 Nanostructured bioceramics

73

Huayna Terraschke, Markus Suta Chapter 5 In situ luminescence analysis of coordination sensors (ILACS): looking inside chemical reactions 105 Huayna Terraschke Chapter 6 In situ monitoring of the syntheses, phase transformations, and loading processes of metal–organic frameworks 133 Huayna Terraschke, Katia Nchimi, Felix Hartmann, Wolfgang Bensch Chapter 7 Operando studies on the charge and discharge processes of battery materials 171 Index

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List of abbreviations AB acac acac ACP ACP ad AFM AMB AOT AP APG ASSB ATR-FTIR BBA BBCDC BCP BDC BDC BESSY BG bipy BIPY bipy BME-BDC bpc BPDC BPEI BPT BPTC bpy BTC BTC BTEC BYPDC BZAC CAM CAU cba ccd CDC CHESS CP CP CPOC CTAB CTAB Cup

4-Aminobenzoate Acetylacetone Acetylacetone Amorphous calcium phosphate Amorphous calcium phosphate Adeninate Atomic force microscopy 4-(Aminomethyl)benzoate Bis(2-ethylhexyl)sulfosuccinate Unsintered apatite Amine-polyglycidol All-solid-state battery Attenuated total reflectance Fourier transform infrared 2-Benzoylbenzoic acid 9,9ʹ-([1,1′-Biphenyl]-4,4ʹ-diyl)bis(9H-carbazole-3,6-dicarboxylate) Bi-phasic calcium phosphate 1,4-Benzenedicarboxylate Benzenedicarboxylate Berliner Elektronenspeicherring Bio-active glass 2,2′-Bipyridine 2,2-Bipyridine 4,4′-Bipyridine 2,5-Bis(2-methoxyethoxy)-1,4-benzenedicarboxylate Bipyridinecarboxylate Biphenyldicarboxylate Branched polyethylenimine Biphenyl-3,4′,5-tricarboxylate 2,2′,5,5′-Biphenyltetracarboxylate Bipyridine 1,3,5-Benzenetricarboxylate Benzenetricarboxylate 1,2,4,5-Benzenetetracarboxylate Bipyridinedicarboxylate Benzoylacetonate Cathode active material Christian-Albrechts-University Carboxyvinyl benzoic acid Charge-coupled device trans-1,4-Cyclohexanedicarboxylate Cornell high-energy synchrotron source Coordination polymer Coordination polymers 5-(4-Carboxyphenoxy)nicotinic acid Cetyltrimethylammonium bromide Cetyltrimethylammonium bromide N-Nitrosophenylhydroxylamine

https://doi.org/10.1515/9783110459098-203

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List of abbreviations

CUS DABCO DBA DB-BDC DCIM DCPD DESY DESY DFLM DFT DHBDC DHBDC dia DMA DMBIM DMF DMF DMF DMNP DOBDC DUT EDTA EDXRD EPFL EPR ESRF EXAFS FE-SEM FITC fu H2BPDC H2BPEB H2FUM H2IDC H2PZDC H3PMBC H4BPTC H4DHTA H6TATPT H6TCPP HAp HCAp HEV HMEIM HPLC HRTEM IBUs ICP-OES ILACS

Coordinatively unsaturated metal sites 1,4-Diazabicyclo[2.2.2]octane 3,4-Dihydroxybenzylamine 2,5-Dibutoxy-1,4-benzenedicarboxylate 4,5-Dichloroimidazolate Dicalcium phosphate dihydrate Deutsches Elektronen-Synchrotron Deutsches Elektronen-Synchrotron Dark-field light microscopy Density functional theory 2,5-Dihydroxybenzene-1,4-dicarboxylate 2,5-Dihydroxyterephthalic acid Diamond topology Dimethylammonium 5,6-Dimethylbenzimidazole Dimethylformamide Dimethylformamide N,N-Dimethylformamide Dimethyl 4-nitrophenylphosphate 2,5-Dioxidobenzene-1,4-dicarboxylate Dresden University of Technology Ethylenediaminetetraacetic acid Energy-dispersive X-ray diffraction École Polytechnique Fédérale de Lausanne Electron paramagnetic resonance European Synchrotron Radiation Facility Extended X-ray absorption fine structure Field emission scanning electron microscope image Fluorescein isothiocyanate Formula unit 4,4′-Benzophenonedicarboxylic acid 1,4-Bis(1H-pyrazol-4-ylethynyl)benzene Fumaric acid 4,5-Imidazoledicarboxylic acid Pyrazoledicarboxylic acid 4-(Phosphonomethyl)benzoic acid 2,2′,5,5′-Biphenyltetracarboxylic acid 2,5-Dihydroxyterdihydroxyterephthalic acid 2,4,6-Tris(2,5-dicarboxylphenylamino)-1,3,5-triazine 4-Tetracarboxyphenylporphyrin Hydroxyapatite Hydroxycarbonate apatite Hybrid electric vehicles 2-Methylimidazole High-performance liquid chromatography High-resolution transmission electron microscope Inorganic building units Inductively coupled plasma emission spectroscopy In situ luminescence analysis of coordination sensors

List of abbreviations

int IPA IR IR IRMOF kat KFSI kgm LAG LAG LB LED LFP LIB Ln lp LPE MB MCPM MeIM MES MIB MIL MIL MIL ML MLMCD MNPs MOF MOF MOF MOF MRI MS MVK NC NCA NCM NDC NGA NIR NMCs NMR NMR NP np OCV ODISC OLEDs

Intermediate Isophthalic acid Infrared Infrared Isoreticular metal-organic framework Katsenite topology Potassium bis(fluorosulfonyl)imide Kagome Liquid-assisted grinding Liquid-assisted grinding Leucomethylene blue Light-emitting diode LiFePO4 Lithium-ion battery Lanthanide Large pore Liquid-phase epitaxy Methylene blue Monocalcium phosphate monohydrate 2-Methylimidazole Mesaconate Magnesium-ion battery Material Institute Lavoisier Material Institute Lavoisier Materials of Institute Lavoisier Metal–ligand Magnetic light-converting molecule device Magnetic nanoparticles Metal-organic framework Metal-organic framework Metal-organic framework Metal-organic framework Magnetic resonance image Mass spectroscopies Methyl vinyl ketone Nanoparticle LiNixCoyAlzO2, where x + y + z = 1 LiNixCoyMnz, where x + y + z = 1 2,6-Naphthalenedicarboxylate Negative gas adsorption Near infrared Nickel manganese cobalt Nuclear magnetic resonance Nuclear magnetic resonance Nanoparticle Narrow pore Open-circuit voltage Oxford-Diamond In Situ Cell Organic light-emitting diodes

IX

X

PCDA PDA PDF PDF PDF PDT PEEK PEG phen PMMA PMOFs PMOs PMT PNC PPY PQDs PTFE PTT PVP pyTa pytz QCM QD SANS SBF SDS SE SEI SEM SEM SERS SHE SIB SNBL SOLEIL sql SURMOF SURMOFs SynRAC TC TCP TEM TEM TEOS THF TM TMB TPA TPU

List of abbreviations

10,12-Pentacosadiynoic acid Calcium 2,6-pyridinedicarboxylate Pair distribution function Pair distribution function Pair distribution function Photodynamic therapy Polyether ether ketone Polyethylene glycol 1,10-Phenanthroline Poly(methyl methacrylate) Porphyrin-based MOFs Periodic mesoporous organosilicas Photomultiplier tube Pre-nucleation clusters Phenylpyridine Perovskite quantum dots Polytetrafluoroethylene Photothermal therapy Polyvinylpyrrolidone Pyridylthioacetate Tridentate pyridine tetrazolate Quartz crystal microbalance Quantum dot Small-angle neutron scattering Simulated body fluid Sodium dodecyl sulfate Solid electrolyte Solid electrolyte interphase Scanning electron microscope Scanning electron microscopy Surface-enhanced Raman spectroscopy Standard hydrogen electrode Sodium-ion battery Swiss Norwegian Beamline Source Optimisée de Lumière d’Energie Intermédiaire du-LURE Square grid Surface-attached metal-organic framework Surface-mounted MOFs Synchrotron-based reaction cell for the analysis of chemical reactions Thioglycolate Tricalcium phosphate Transmission electron microscope Transmission electron microscopy Tetraethyl orthosilicate Tetrahydrofuran Transition metal Trimethylbenzene Terephthalic acid Thermoplastic polyurethane

List of abbreviations

TRLFS tta TTA UiO UiO UiO UTSA WLED XANES XAS XAS XRD XRD XRD XRD XRD ZBFB ZIF ZIF

Time-resolved laser-induced fluorescence spectroscopy 2-Thenoyltrifluoroacetylacetonate Thenoyltrifluoroacetonate University of Oslo University of Oslo University of Oslo University of Texas at San Antonio White-light-emitting diode X-ray absorption near-edge spectroscopy X-ray absorption X-ray absorption spectroscopy X-ray diffraction X-ray diffraction X-ray diffraction X-ray diffraction X-ray diffraction Zinc-bromide redox-flow batteries Zeolitic imidazolate framework Zeolitic imidazolate framework

XI

List of contributors Huayna Terraschke, Jun.-Prof. Dr. Institute of Inorganic Chemistry Christian-Albrechts-Universität zu Kiel Max-Eyth-Straße 2, 24118 Kiel Germany E-Mail: [email protected] Katia Nchimi Nono, Dr. Institute of Inorganic Chemistry Christian-Albrechts-Universität zu Kiel Max-Eyth-Straße 2, 24118 Kiel Germany and Department of Inorganic Chemistry Faculty of Science, University of Yaoundé I P.O. Box 812 Yaoundé Cameroon Moritz Bassen, B.Sc. Institute of Inorganic Chemistry Christian-Albrechts-Universität zu Kiel Max-Eyth-Straße 2, 24118 Kiel Germany Patric Lindenberg Institute of Inorganic Chemistry Christian-Albrechts-Universität zu Kiel Max-Eyth-Straße 2, 24118 Kiel Germany Niklas Ruser, M.Sc. Institute of Inorganic Chemistry Christian-Albrechts-Universität zu Kiel Max-Eyth-Straße 2, 24118 Kiel Germany

https://doi.org/10.1515/9783110459098-204

Markus Suta, Jun.-Prof. Dr. Inorganic Photoactive Materials Institute of Inorganic Chemistry and Structural Chemistry Heinrich Heine University Düsseldorf Universitätsstraße 1, 40225 Düsseldorf Germany Felix Hartmann, Dr. Institute of Inorganic Chemistry Christian-Albrechts-Universität zu Kiel Max-Eyth-Straße 2, 24118 Kiel Germany and Physikalisch-Chemisches Institut & Zentrum für Materialforschung (ZfM) Justus-Liebig-Universität Gießen Heinrich-Buff-Ring 17, 35392 Gießen Germany Wolfgang Bensch, Prof. Dr. Institute of Inorganic Chemistry Christian-Albrechts-Universität zu Kiel Max-Eyth-Straße 2, 24118 Kiel Germany

Katia Nchimi, Huayna Terraschke

Chapter 1 Introduction In general, nanomaterials have been gaining great attention in the past few years since they can be used in various technological and medical applications [1]. For example, nanostructured materials are applied on the technological development of energy-saving devices (e.g. light-emitting diodes (LEDs)), renewable energy sources (e.g. solar cells), energy storage (e.g. rechargeable batteries), and gas storage solutions (e.g. hydrogen fuel cells), which are crucially important to reduce the impact of global warming and overcome energy shortage issues. Furthermore, nanomaterials have been widely used in medicine, particularly in the fields of optical imaging and drug delivery. Nanostructured materials have at least one dimension that is smaller than 100 nm, and include various examples such as nanoparticles (NPs), thin films, or metal-organic frameworks (MOFs) with nanosized voids [2, 3]. NPs have a relatively small number of constituent atoms and a high surface area. Consequently, any subtle addition, removal, or modification of a few atoms, as for example, through size and shape alteration can have significant effects on the interaction with e.g. electromagnetic radiation and can lead to profound changes on the material’s properties [4]. Among the nanostructured materials, MOFs are extremely interesting because they can also be functionalised and tailored e.g. via the adsorption of other chemical compounds within their pores. The nanomaterial structure-related properties can be regulated by, for instance, controlling the synthesis conditions. Their precise syntheses require profound knowledge and strict control over the respective crystallisation processes, which can be achieved by developing techniques and reactors to in situ monitor the respective chemical reactions [5]. The mechanism of chemical reactions could be studied by removing samples during the process from the reactor and analysing them ex situ, which is rather disadvantageous as it only offers snapshots of the process with very limited time resolution and can influence the sample upon its preparation for the analysis. These issues can be overcome by performing in situ measurements during the reactions. In general, in situ characterisation techniques are essential for elucidating the different phenomena behind the product formation, such as nucleation, crystal growth, the formation of (amorphous) intermediates, or polymorphic phase transformations, thus also allowing for the detection of new metastable compounds and for the optimisation of the synthesis conditions [5]. This book is directed to undergraduate and graduate students, aiming to obtain an overview on the research field of inorganic nanostructured materials. The book is divided into two parts: Part I (Chapters 2–4) offers a summary of the synthesis, properties, and applications of various classes of nanostructured compounds, while Part II https://doi.org/10.1515/9783110459098-001

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Katia Nchimi, Huayna Terraschke

(Chapters 5–7) reviews the applications of in situ analysis elucidating, for example, formation, transformation and operation of such substances. In detail, Chapter 2 focuses on the guest-induced luminescence of MOFs, which is typically used in lighting and chemical sensing applications. The introduction of luminescent materials, such as lanthanide ions or complexes, quantum dots, plasmonic nanoparticles, and organic phosphors into the MOF pores increases the quantum efficiency and chemical stability, and results in fine-tuning of the emission wavelength [6]. The remaining chapters in Part I explore the scope of nanostructured materials beyond their optical properties, focusing on their magnetism and biocompatibility. Chapter 3 summarises for instance the strategies for the synthesis and functionalisation, as well as the applications of one of the most used types of magnetic nanomaterials, the iron oxide-based NPs. These NPs are of utmost importance in the biomedical and technological fields, such as magnetic resonance imaging, cancer treatment, drug delivery, bioimaging, and magnetic separation, in addition to their application as energy materials and nanocatalysts [7]. Chapter 4 describes the different aspects of the synthesis, functionalisation, and applications of nanostructured bioceramics. These particles comprise calcium phosphates and carbonates, glasses, or porous materials, and their high biocompatibility allows for their application not only in bioimaging and drug delivery but also in the production of bone cements and coatings for orthopaedic and dental implants [8]. In Part II, Chapter 5 presents the opportunity to apply in situ luminescence measurements for monitoring of the formation of functional materials ranging from complexes to bioceramics. One important approach in this field is the in situ luminescence analysis of coordination sensors (ILACS), which explores the sensitivity of the optical properties of lanthanide ions for implanting them as local crystallisation tracer during the formation of solid matter. This approach is particularly advantageous because of its high temporal resolution, sensitivity, simplicity, and high accessibility [9]. Chapter 6 reviews the available in situ techniques and applications for tracking the crystallisation of MOFs, as well as their structural transformations upon the loading of guest molecules. The main techniques discussed include in situ X-ray diffraction, X-ray absorption, and total scattering, in which high-energy and high-flux synchrotron radiation allows for the permeation of X-rays through the reactor walls and bulk solvents [10]. Finally, Chapter 7 offers an overview of examples of in situ analytical tools applied during battery operation, also known as the operando characterisation techniques. Performing operando studies during the charge and discharge processes enables a deep understanding of the working principle of batteries and the reason behind their failure. Such measurements are important for promoting electromobility and storing power from fluctuating renewable sources such as solar and wind energy [11].

Chapter 1 Introduction

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References [1]

Z. Wu, S. Yang, W. Wu, Shape control of inorganic nanoparticles from solution. Nanoscale 8, 1237–1259, 2016 and references therein. [2] M. Auffan, J. Rose, J.-Y. Bottero, G. V. Lowry, J.-P. Jolivet, M. R. Wiesner, Towards a definition of inorganic nanoparticles from an environmental, health and safety perspective. Nat. Nanotech. 4, 634–641, 2009. [3] O. M. Yaghi, M. O’Keeffe, N. W. Ockwig, H. K. Chae, M. Eddaoudi, J. Kim, Reticular synthesis and the design of new materials. Nature 423, 705–714, 2003. [4] F. P. G. de Arquer, D. V. Talapin, V. I. Klimov, Y. Arakawa, M. Bayer, E. H. Sargent, Semiconductor quantum dots: Technological progress and future challenges. Science 373, 640/1–15, 2021. [5] N. Pienack, W. Bensch, In-Situ Monitoring of the Formation of Crystalline Solids. Angew Chem. Int. Ed. 50, 2014–2034, 2011. [6] G. H. Carey, A. L. Abdelhady, Z. Ning, S. M. Thon, O. M. Bakr, E. H. Sargent, Colloidal Quantum Dot Solar Cells Chem. Rev. 115, 12732–12763, 2015. [7] S. Wang, J. Xu, W. Li, S. Sun, S. Gao, Y. Hou, Magnetic nanostructures: rational design and fabrication strategies toward diverse applications. Chem. Rev. 122, 5411–5475, 2022. [8] G. L. Koons, M. Diba, A. G. Mikos, Materials design for bone-tissue engineering. Nat. Rev. Mater. 5, 584–603, 2020. [9] H. Terraschke, M. Rothe, P. Lindenberg, In-Situ Monitoring Metal-Ligand Exchange Processes by Optical Spectroscopy and X-Ray Diffraction Analysis: A Review. Rev. Anal. Chem. 37, 20170003/1-22, 2018. [10] H. Reinsch, N. Stock, Synthesis of MOFs: a personal view on rationalisation, application and exploration. Dalton Trans. 46, 8339–8349, 2017. [11] E. Pomerantseva, F. Bonaccorso, X. Feng, Y. Cui, Y. Gogotsi, Energy storage: The future enabled by nanomaterials. Science 366, 969/1-12, 2019.

Moritz Bassen, Katia Nchimi, Huayna Terraschke✶

Chapter 2 Guest-induced luminescence of metal-organic frameworks The loading of guest molecules into the porous structures of metal-organic frameworks (MOFs) permits the adjustment and improvement of their luminescence properties, in particular, their quantum yield and precise tuning of the emission wavelength. Despite the importance of this research field and its promising potential applications, to the best of our knowledge, only few review articles have been published to date on guestinduced MOF luminescence, partially limited to sensor-related applications. This chapter offers therefore an extensive survey of the existing works in this research area, discussing its current challenges, aiming to substantially contribute to its further development.

2.1 Introduction Over the past few decades, there has been a constant increasing interest in metal-organic frameworks (MOFs). These materials are made of metal ions or clusters connected to each other by organic linkers via coordinative bonds to create a defined open crystalline framework with permanent porosity [1]. MOFs exhibit ultra-high porosity (up to 90% free volume) and enormous internal surface areas, extending beyond 6,000 m2/g [2], exceeding those of traditional porous materials, such as carbon-based ones and zeolites. MOFs with permanent porosity exhibit more variety and multiplicity than any other class of porous materials because they can be composed of several different ligands containing different types and amounts of functional groups [1]. MOF materials exhibiting a large variety of sizes and 3D framework structures can be constructed based on the ligands and the central metal ions used. This versatility makes them advantageous for many applications, such as gas storage, separation, catalysis, chemical sensing, drug delivery, and other biomedical applications. Therefore, MOFs have been extensively studied and >20,000 different structures have been reported to date [1]. In general, MOFs are interesting because of their intrinsic ability to emit light, either ligand-based (light is emitted by conjugated π-systems within the organic ligands [3, 4]) or metal-based (light is emitted through the metal centres within the MOF structure [5–7]). A third possible emission derived from the porous structure of MOFs, which allows guests to be inserted, is the so-called guest-induced luminescence. ✶

Corresponding author: Huayna Terraschke, E-Mail: [email protected]

https://doi.org/10.1515/9783110459098-002

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

The guest can show luminescence directly or can lead to complex energy transfer processes inside the MOFs, thereby changing the optical properties of the host structure. These features make MOFs excellent candidates for additional applications based on luminescence, such as light-emitting devices [8]. Upon comparison, emission from single organic or single inorganic luminescent materials often covers only a narrow part of the visible light spectrum, while MOFs offer relatively high luminescence intensities and broad emission spectra through the association of inorganic metals with organic linkers. The combination of well-defined porosity and environmentally dependent optical properties reinforces the interest in MOFs as versatile optically active materials [8]. Despite all of the above-mentioned advantages, so far there are, to the best of our knowledge, only few review articles on guest-induced luminescence. One important example was published by Allendorf et al. [9] and another was recently published by Shu et al. [10]. However, they mainly focus on their chemical sensing applications. Therefore, the objective of our present work is to offer a general overview of the recent research on guest-induced luminescence in MOFs, focusing on lanthanide ions, lanthanide complexes, quantum dots, plasmonic nanoparticles, and organic phosphors as guests inside the pores of MOFs. The synthesis and loading methods are first presented before discussing the most commonly reported materials in this field. Finally, the recent applications and future research possibilities are discussed.

2.2 Synthesis and loading methods 2.2.1 MOF synthesis As discussed in detail in Chapter 6, MOFs consist of two major components: metal ion centres and bridging organic linkers. MOFs are prepared by combining metal ions with organic linkers, mostly under mild reaction conditions. According to Soni et al. [11], the main synthesis methods can be categorised into conventional (solvothermal) and unconventional methods (see Figure 2.1). In general, using conventional solvothermal synthesis, a mixture of solvated metal ions and organic linkers is heated in a glass vial, Teflonlined autoclave, or pressure reactor [12]. If water is used as the solvent, this method is known as hydrothermal synthesis. For a specific structure, controlling the synthesis parameters, such as pressure, temperature, and solvent composition, is important. On the other hand, mechanochemical synthesis is an unconventional method. It works by breaking the intramolecular bonds using mechanical forces, followed by chemical transformation. A mixture of the metal salt and organic linker is ground in a mortar and pestle or a ball milling device without the use of solvent. Next, the mixture is gently heated to evaporate any water or other volatile molecules that are formed as

Chapter 2 Guest-induced luminescence of metal-organic frameworks

7

by-products. If solvents are added during the grinding step, it is known as liquidassisted grinding [11]. Further synthetic routes have been developed, in addition to these two main methods. Four important examples are microwave-assisted, electrochemical synthesis, sonochemical, and layer-by-layer syntheses [11]. Microwave-assisted synthesis uses energy provided in the form of microwave radiation, allowing fast crystallisation, narrow particle size distribution [13], and morphological control [14, 15]. The electrochemical approach uses metal ions, which are continuously supplied using anodic dissolution, instead of metal salts [16]. These ions react with the dissolved linker molecules and the conducting salt in the reaction medium. Sonochemical synthesis is a rapid and environmentally friendly method in which ultrasonic radiation is used. Using homogenous and accelerated nucleation, a reduction in the crystallisation time and significantly smaller particle sizes can be achieved than those obtained using conventional solvothermal synthesis.

Figure 2.1: Overview of the synthesis methods, possible reaction temperatures, and the final reaction products in the synthesis of MOF [17] (Copyright 2012 American Chemical Society).

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

Finally, a layer-by-layer method has been used to form thin MOF films. It is based on surface chemistry, in which a functionalised organic surface is immersed sequentially into solutions of metal ions and organic linkers [11].

2.2.2 Encapsulation of guest species There are three main approaches used to construct MOF composites containing guest species in their pores. These are known as “ship in a bottle”, “bottle around ship”, and “in situ one-step synthesis” (Figure 2.2). In addition, guest molecules and/or ions can diffuse into MOFs from the solution or gaseous phase. The “ship in a bottle” approach involves the encapsulation of guest species by first loading the cavities of an existing host MOF with small molecule components or nanoparticle precursors. Subsequently, further treatment is performed, leading to the desired functional structure of the encapsulated guest species. Therefore, for NP–MOF guest–host structures, the nanoparticles (ships) are assembled inside the MOF (bottle). Various techniques, such as solution infiltration, solid grinding, and vapour deposition, have been used to introduce the NP precursors into MOFs [18, 19]. This loading approach is challenging because it is difficult to precisely control the location, composition, structure, and morphology of the incorporated guests. However, encapsulation in the MOF pores, which limits the growth of the molecules and nanoparticles inside, can be advantageous, especially during the synthesis of size-dependent guest species, such as quantum dots [20]. The “bottle around ship” approach, which is also known as template-assisted synthesis, involves two steps. First, functional molecules or nanoparticles are prepared and often stabilised using capping agents or surfactants. In the second step, the presynthesised guest species are added to a synthetic solution containing the MOF precursors used to assemble the MOF. The guest species do not occupy the cavities of the MOF, but are rather surrounded by grown MOFs. The MOF (bottle) is built around the guest species (ship). With this method, the problem of nanoparticle aggregation on the external surface of the MOF is limited. Furthermore, the size, morphology and structure of the entrapped nanoparticles can be easily controlled because they are synthesised prior to the assembly of the MOF framework [19, 20]. On the other hand, the introduction of nanoparticles can result in difficulties during the subsequent growth of the MOFs because of the interfacial energy barriers between the two types of materials. In addition, the presence of capping agents may be unfavourable for the complete exposure of the active sites and may decrease the performance of the nanoparticles [20]. The third synthesis method used is the in situ one-step approach. This involves the direct mixing of the precursor solutions that comprise the guest species and MOFs. The growth and assembly of the nanoparticles and MOF occur simultaneously. When compared to the previously mentioned synthetic approaches, this strategy is

Chapter 2 Guest-induced luminescence of metal-organic frameworks

9

Figure 2.2: The main approaches used for the fabrication of nanoparticle/MOF composites: (a) ship in a bottle, (b) bottle around ship, and (c) in situ one-step synthesis (reprinted with permission from reference [20]).

straightforward and simple, but requires an appropriate balance between the particle growth and self-nucleation of the nanoparticles and MOF. Therefore, the choice of the functional group in the organic linkers or solvent is crucial for trapping the nanoparticle/guest species precursors, stabilising the species formed in situ, and facilitating the hetero-nucleation of the MOFs on the surface of the guest species [21, 20].

2.3 Materials used for guest-induced MOF luminescence 2.3.1 MOF encapsulation of lanthanide ions Lanthanides are a class of elements in the periodic table from atomic number 58 (cerium) to 71 (lutetium). Complemented by scandium, yttrium, and lanthanum, they form a group of rare-earth elements. Lanthanum shares the same line as the 14 lanthanides in the periodic table, but it does not belong to the same family because its 4f orbital is empty, while from Ce to Lu, the 4f orbital is partially or fully filled. Through ionisation to the most common oxidation state (+3), a uniform configuration [Xe]4f n−1 is achieved in this series. The partially filled valence orbitals of the trivalent ions from Ce3+ to Yb3+ lead to the possibility of 4f → 4f or 4f → 5d electronic transitions, resulting in luminescence in the UV to IR region. The luminescent properties of lanthanide ions that will be discussed in the remainder of this chapter are focused on the

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

4f–4f transitions. Each electronic configuration corresponds to an energy state, and the Dieke diagram (Figure 2.3) shows the ground and excited states of the trivalent lanthanide ions. For instance, it is used to assign the emission peak to its respective electronic transition [22]. The 4f orbitals are the valence orbitals that are shielded by the filled 5s and 5p outer orbitals. Therefore, lanthanides bind mostly through ionic interactions, and the influence of the ligand field on the 4f orbitals is minimal. For this reason, a low variability of the 4f–4f emission properties of trivalent lanthanides is expected, which allows these metal ions to have unique luminescent properties, such as narrow emission peaks, large Stokes shifts, strong emission intensities, and long lifetimes. These optical properties have motivated an increasing amount of interest in luminescent lanthanide compounds. Therefore, apart from the extensive research on MOFs constructed using lanthanide ions as the metal centres within the building units, an increasing number of reports have been published on the application of lanthanide ions as guest species doped in the MOF structure. The encapsulation of lanthanide ions is especially advantageous because the luminescence of pure lanthanide ions is severely quenched when exposed to an aqueous environment. They are effectively shielded against such effects when they are encapsulated inside the pores of the MOF [10]. In addition, pure lanthanide ions show relatively low quantum yields because they suffer from a weak absorbance due to the parityforbidden electronic transitions. A common way to circumvent this problem is to couple lanthanide ions with the strongly absorbing linkers of MOFs, allowing direct energy transfer from the more readily accessible linker excited states to the appropriate lanthanide metal energy level. In this way, specific and narrow emission peaks with high quantum yields can be achieved [24]. For example, europium and terbium trivalent ions are the most common lanthanide ions reported in guest–MOF composites because of their strong emission in the visible light region. MOFs derived from aromatic cores, such as benzene, biphenyl, or bipyridine, and connected with other metal ions are generally efficient for the generation of guestinduced luminescence with lanthanide ions. For example, an anionic microporous MOF based on the association of 1,2,4,5-benzenetetracarboxylate (BTEC) with Zn2+ has been prepared and loaded with various lanthanides by Luo and Batten [25], which demonstrated the ability to sensitise the ligands on both Eu3+ and Tb3+. These lanthanide–MOF composites were obtained from the hydrothermal self-assembly of Zn(NO3)2 · 6H2O, pyromellitic dianhydride, and (NH4)6Mo7O24 · 4H2O, followed by cation exchange of the NH4+ in the structure by lanthanide ions. The authors also underlined the tuneable properties of such materials from the original blue emission to red or yellow emission, depending on the choice and concentration of the dopant used (Eu3+ or Tb3+) [25]. The well-known Bpy-UiO MOF based on the polymerisation of a bipyridinecarboxylate (bpc) linker with Zr4+ was reported by Zhou et al. [26]. It acts as a responsive luminescent composite upon the incorporation of Eu3+ cations into its porous nanostructure. In this guest–host assembly, strong coordination interactions formed between Eu3+ and the free

Chapter 2 Guest-induced luminescence of metal-organic frameworks

11

Figure 2.3: Dieke diagram of the energy levels of trivalent lanthanide ions (republished with permission of Royal Society of Chemistry from reference [23]; permission conveyed through Copyright Clearance Center, Inc.). On the abscissa, the different trivalent lanthanide ions are entered. On the other hand, the ordinate shows the energy, displayed as the wavenumber of the entered energy states of the ions. The red-marked energy states are those from which a radiative emission to the ground state is most likely to be observed.

bipyridyl moieties construct a stable structure along with efficient sensitisation, and hence good luminescence properties centred on europium either in the solution or solid state. The loading of another metal ion simultaneously with the lanthanide in a given MOF is also possible and, in general, permits the tuning of the emission of the material. Hao et al. [27] prepared a UiO-66(Zr)-(COOH)2 MOF using the bpc ligand, loading it with Ag+ and Eu3+ ions. The host MOF was prepared using a sol–gel process, upon heating a mixture of bpc and ZrCl3 at reflux for 24 h in an air atmosphere. After centrifugal separation, the materials were dried and co-doped with Ag+ and Eu3+ in various amounts, upon immersion of the MOF in a solution of the nitrate salts of the 4d and 4f ions. When Eu3+ ions were doped alone, the photoluminescence studies indicated that an efficient ligand-to-Eu energy transfer occurred because the emission spectra exhibited the extinction of the residual blue emission of the ligand along with the intense and characteristic sharp emission peaks of the encapsulated Eu3+ ions. In contrast, the bimetallic Ag+/Eu3+-UiO-66(Zr)-(COOH)2 guest MOF exhibited the appearance of ligand fluorescence, while the emission of europium was drastically reduced [27]. Another interesting bimetallic ion guest has been described by An et al. [28], who studied the encapsulation of trivalent lanthanide ions into the bio-MOF-1 structure to achieve a tuneable MOF luminescence by combining different encapsulated Ln3+ ions. In this work, they prepared a series of Ln3+@bio-MOF1 (Ln3+ = Tb3+, Sm3+, Eu3+, or Yb3+) materials using a cation exchange process. For this purpose, the dimethylammonium cations in the one-dimensional channels of bio-MOF-1 ([Zn8(ad)4(BPDC)6O(2 Me2NH2) · 8DMF · 11H2O]; ad, adeninate; BPDC, biphenyldicarboxylate) were exchanged with lanthanide ions by soaking the as-synthesised bio-MOF-1 samples in a DMF solution of Ln(NO3)3. The

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

doped MOFs showed characteristic sharp emissions, corresponding to their respective encapsulated lanthanide cations, when excited at 340 nm. Tb3+ exhibited its typical green emission at 545 nm, Sm3+ exhibited an orange-pink emission at 640 nm, Eu3+ a red emission at 614 nm, and Yb3+ at 970 nm (NIR) (Figure 2.4). Their emission colour could be observed by the naked eye during excitation under a standard laboratory UV lamp (365 nm). The similarity of all four lanthanide-centred excitation spectra, with a consistent maximum at 340 nm, suggested that the energy migrated through the same electronic levels located in the MOF structure. These characteristic emission peaks were also easily detected even in an aqueous environment with considerably high quantum yield, even though water molecules normally exhibit a strong quenching effect on the luminescence of non-encapsulated lanthanide ions. This work demonstrated that the MOF, as explained above, does not only serve as an antenna for sensitising three different visible-emitting and one NIR-emitting lanthanide cations, but can also protect lanthanide cations from solvent quenching. This proves that the encapsulation of lanthanide ions into the MOF pores is highly advantageous [29]. Additional examples of lanthanide ions, doped or encapsulated within the porous MOF structure, which exhibit guest-induced luminescence properties are presented in Table 2.1.

Figure 2.4: Colour-tuneable luminescence of lanthanide ions encapsulated by bio-MOF1 [30] (Copyright 2011 American Chemical Society).

2.3.2 MOF encapsulation of metal ion complexes In addition to free metal ions, coordination complexes of various metallic species, including lanthanides, transition metals, and post-transition metals such as gallium or aluminium, can be introduced or associated with ligands as guest compounds in MOFs to exhibit guest-induced luminescence. In general, the selected complexes (Table 2.2) already

Chapter 2 Guest-induced luminescence of metal-organic frameworks

13

Table 2.1: A summary of lanthanide ions doped in MOF compounds. Lanthanide ion

Metal centre

Linker

Ln+ = Tb+, Sm+, Bio-MOF- Eu+, or Yb+

Zn+

Biphenyldicarboxylate (BPDC)

[]

Eu+

Al+

Benzenedicarboxylate (BDC)

[]

Benzenetricarboxylate (BTC)

[]

Eu

Host MOF

MIL--COOH(Al)

+

MIL-

+

+

Ga

+

References

UiO-

Zr

,-Pyridinedicarboxylate

[]

Eu+

Uio--(COOH)

Zr+

,,,-Benzenetetracarboxylate (BTEC)

[]

Eu+

Bpy-UiO

Zr+

Bipyridinedicarboxylate (BYPDC)

[]

Eu+

UiO-(Zr)-(COOH)

Zr+

BYPDC

[]

BYPDC

[]

BYPDC

[]

BYPDC

[]

Eu

Eu Eu Tb

+

UiO-bpydc

+

UiO--bpydc

+

Al(OH)-(bpydc)

+

+

Zr

+

Zr

+

Al

+

Zn-MOF

Zn

,′,′′-[(,,-Triazine-,,-triyl)tris(sulfanediyl)]tribenzoic acid

[]

Tb+

Zr-MOF

Zr+

Terephthalic acid and isophthalic acid

[]

Tb+

[Cu(HCPOC)]n

Cu+

-(-Carboxyphenoxy)nicotinic acid) (CPOC)

[]

Eu+/Tb+

Ln(OH)(BYPDC)

In+

BYPDC

[]

Tb

+

+

+

Eu /Tb

Bio-MOF-

Zn

BYPDC

[]

Eu+/Tb+

Zn(II)-MOF

Zn+

Dimethyl--hydroxy-benzene-,dioate

[]

possess emission properties, which contribute to modifying the emission features of a given MOF. For example, Yin et al. [45] worked on a composite comprised of Ru(bpy)32+ (tris (2,2bipyridyl)ruthenium(II)), doped into the pores of MIL-101(Al)-NH2 and used for the ratiometric fluorescence sensing of water in various organic solvents. While the ruthenium complex emits red light at 615 nm, MIL-101(Al)-NH2 exhibits a light blue emission at 465 nm. In the synthesis, tris(2,2-bipyridyl)dichloride ruthenium(II)hexahydrate [Ru(bpy)3Cl2 · 6H2O] was added to a solution of aluminium chloride hexahydrate (AlCl3 · 6H2O) and 2-aminoterephthalic acid (BDC−NH2) in DMF. The mixture was transferred to a Teflon-lined autoclave, sonicated for 5 min, and maintained at 130 °C for 72 h. Later, the solvent was removed via centrifugation and washed with DMF and ethanol. The

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

Figure 2.5: A schematic illustration of the enhanced blue emission observed in the presence of water using Ru@MIL-101(Al)-NH2 [45] (Copyright 2017 American Chemical Society).

Figure 2.6: A schematic representation of the formation of hierarchically structured hybrid MOF particles encapsulating Tb complexes (Tb complex@ZIF-8@TIF-1Zn-PDA) (republished with permission of Royal Society of Chemistry from reference [46]; permission conveyed through Copyright Clearance Center, Inc.).

MOF prepared via this one-step in situ method showed dual emission. However, the intensity of the blue fluorescence was enhanced upon increasing the water content in the organic solvents studied. This is due to the sensitivity of BDC−NH2, upon protonation in the presence of water. On the other hand, the red fluorescence of Ru(bpy)32+ remained

Chapter 2 Guest-induced luminescence of metal-organic frameworks

15

stable. In this way, turn-on sensing of the water content was achieved using the ratiometric fluorescence of the dual emission under a single excitation of 300 nm utilising Ru@MIL−NH2 as the probe (Figure 2.5) [45]. Another interesting approach for the encapsulation of complexes by an MOF was reported by Wang et al. They encapsulated [Tb(pytz)3] (pytz, tridentate pyridine-tetrazolate) complexes into the framework of zeolitic imidazolate framework-8 (ZIF-8) nanoparticles. Interestingly, the ZIF-8 nanoparticles were encapsulated into the TIF-1 Zn MOFs via complexation. The [Tb(pytz)3] complex was encapsulated using an in situ method, utilising Zn2+ and 2-methylimidazole (HMIM), to form the ZIF-8 MOF. Subsequently, the hydrophobic ligand, 5,6-dimethylbenzimidazole (DMBIM), was used to induce the aggregation of the ZIF-8 nanoparticles through the terbium complexes absorbed on their surfaces and, in turn, the hybridisation of the TIF-1Zn (Zn(DMBIM)2) framework (Figure 2.6). While the encapsulation of the Tb complex in ZIF-8 was not sufficient to protect the complex from environmental species, such as water in biological media, a second encapsulation in the TIF-1Zn framework shows sufficient protection because the bridging ligands in TIF1Zn are hydrophobic [46]. More examples of how the luminescence of guest–MOF compounds can be used for chemical sensing applications will be described in the applications section of this chapter.

2.3.3 MOF encapsulation of quantum dots Semiconductor nanocrystals with sizes ranging from 2 to 10 nm are often referred to as quantum dots (QDs). They have received significant attention due to their tuneable and size-dependent electronic properties. They are often aggregates of hundreds and thousands of atoms, mainly from groups II to VI of the periodic table, such as CdSe, ZnSe, or ZnO, groups III to V (such as InP and InAs), or from groups IV to VI (such as PbSe) [55]. The optical properties of QDs are strongly influenced by their size. While a CdSe nanocrystal with a diameter of 2 nm shows blue fluorescence, a larger crystal of ~ 6 nm emits red light. This effect is based on the fact that different particle sizes lead to a change in the band gap energy. Smaller particles have relatively large band gaps and therefore show high energetic luminescent emission in the blue/UV range. Larger particles show red/NIR emission because the band gap is smaller and less energetic (Figure 2.7). The number of atoms in the QD particles increases with the diameter of the particle, increasing the number of different energetic atomic orbitals in these materials. The energy band increases and the band gap decreases, resulting in a red shift [56]. Due to the above-mentioned advantages, QDs are of significant interest as possible guests to fill the pores of MOFs in order to improve and adjust their luminescent properties. Here, the most important examples of group II–VI semiconductors, Group

16

Table 2.2: A summary of complexes encapsulated in MOFs. Host MOF

Metal centre

Linker

[Eu(BZAC)BPY]

HKUST-

Cu+

[Gallium(III) tris(-hydroxyquinolinato)] (Gaq)

ZIF-

References

Benzene-,,-tricarboxylic acid (HBTC)

[]

+

-Methylimidazole

[]

+

,,-Tris (,dicarboxylphenylamino)-,,-triazine (HTATPT)

[]

Zn

[Ir(Hppy)(bpy)] (Hppy, -phenylpyridine; bpy, ,′-bipyridine)

[(CH)NH][(CdCl)(TATPT)] · DMF · HO

Cd

Chloroacetic acid (Ln = Eu, Nd)

IRMOF- → Ln-IRMOF--CA

Zn+

-Aminoterephthalic acid

[]

+

-Aminoterephthalic acid

[]

+

-Aminoterephthalic acid

[]

+

-Aminoterephthalic acid

[]

,,,-Benzenetetracarboxylic acid (BTEC)

[]

+

Glyoxylic acid (Ln = Eu, Nd) Diethyl (ethoxymethylene)malonate (Ln = Eu, Nd) Methyl vinyl ketone (Ln = Eu, Nd)

Ln-IRMOF--GL (Ln = Eu, Nd) Ln-IRMOF--EM (Ln = Eu, Nd) Ln-IRMOF--MVK (Ln = Eu, Nd)

Zn Zn Zn

+

Tb(pytz) (pytz, tridentate pyridine-tetrazolate)

ZIF-

Al

PDA-Eu-BBA-PEG-DBA-FeO (PDA, calcium-,pyridinedicarboxylate; BBA, -benzoylbenzoic acid; PEG, polyethylene glycol; DBA, ,-dihydroxybenzylamine)

ZIF-

Zn+

-Methylimidazole

[]

[Ru(bpy)]+ (bpy, bipyridine)

UiO-

Zr+

, ′-Bipyridine-,-dicarboxylic acid (BPDC)

[]

Moritz Bassen, Katia Nchimi, Huayna Terraschke

Complex composition

[Ru(bpy)]+ (bpy, bipyridine)

MIL--(Al)-NH

Al+

,-Terephthalic acid (HBDC)

[]

[Ru(bpy)]+ (bpy, bipyridine)

UiO-

Zr+

HBDC

[]

+

[Ir(CF-PPY-F)(bpy)] (PPY, phenylpyridine; bpy, bipyridine)

(MeNH) [Zn(L) (HO)] · DMA

Zn

,-(-(-Carboxyphenylamino)-,,triazine-,-diyl-diimino) diterephthalic acid (HL)

[]

Tetraamine platinum hydroxide

HMIL-

Al+

BTEC

[]

+

Chapter 2 Guest-induced luminescence of metal-organic frameworks

17

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

IV (carbon) QDs, and halide perovskite composites are explained in detail. Among the different types of semiconductor nanocrystals, chalcogenide-based QDs, generally with sulphur (S), selenium (Se), and tellurium (Te), are the most promising and studied materials, not only because of their potential applications in optoelectronics and biomedicine, but also because of their well-known chemistry and established protocols used for their preparation, which allow meticulous control of their shape, size, and structure [55]. Moreover, the encapsulation of various QDs into the MOF matrix can result not only in a tremendous increase in their photoluminescence, but also in the alleviation of the aggregation-caused quenching of the luminescence that is observed with free QDs [58]. One interesting example of a QD-loaded compound is QD@MOF, which was developed by Jin et al. [59] The aim of their work was to develop a QD@MOF suitable for light harvesting, in which solar energy is converted into electrical or chemical energy. A promising class of materials in this regard is metalloporphyrin MOFs, whose porphyrin-derived linkers are structurally similar to naturally occurring pigments, such as chlorophyll and haemoglobin, and light harvesting moieties. The efficiency of these components depends on the number of photons absorbed from the solar spectrum. As metalloporphyrin provides only a limited coverage of this spectrum due to its narrow absorption band, their combination with semiconductor NPs with narrow band gaps and significantly broader absorption spectra can offer several advantages. Therefore, two Zn/porphyrin-based MOFs were functionalised with amine-capped CdSe/ZnS core–shell QDs. In this case, the QD emission band overlaps with the MOF absorption band, indicating that resonant energy transfer from the QD donor to the MOF acceptor is possible. In addition, it is important to note that the MOF component of the QD@MOF composites emitted when excited at wavelengths where the MOF composite alone does not absorb, confirming the light harvesting and energy transfer effects. Energy transfer times between 1.7 and 10 ns were observed, with efficiencies as high as 84% being reached. Finally, this work showed that the light harvesting efficiency can be tuned by modulating the MOF composition or QD size, both of which influence the spectral overlap [59]. Carbon QDs are small particles, typically with diameters of 5,000 K (daylight) [62, 117].

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

Table 2.4: A summary of the organic phosphors encapsulated in MOF composites reported up until 2021. Organic phosphor

Host MOF

Metal Linker centre

,-Bis(-methyl-(-formylphenyl)thiophen-yl) hexafluorocyclopentene (OF-DAE)

ZJU-

Eu+

,′:′,″:″,‴-Quaterphenyl -,‴,,‴-tetracarboxylicacid (QPTCA)

[]

-(pDimethylaminostyryl) --methylpyridinium (DSM) + acriflavin

ZJU-

In+

,′,″-Benzene-,,-triyl-tribenzoate (HBTB)

[]

-[p-(Dimethylamino) In-MOF styryl]--ethylpyridinium (DSM)

In+

,′,″-Nitrilotribenzoic acid (HTNB)

[]

-Aminonaphthalimide

LnL

Ln = Eu+, Gd+

–

[]

-Aminofluorescein

MOF-

Zr+

Fumaric acid

[]

Fumaric acid

[]

Calcein

MOF-

+

Zr

+

References

Coumarin 

Ln-BPT

Eu

Biphenyl-,′,-tricarboxylate (BPT)

[]

Coumarin 

Ln-BPT

Tb+

BPT

[]

Coumarin 

Ln-BPT

EuxTby BPT

[]

Eosin Y Fluorescein

DUT- EuL

Fluorescein + rhodamine ZIF-

+

Zr

Eu

+

Zn+ +

,-Naphthalenedicarboxylic acid [,′-(Carbonylbis (azanediyl))bis(methoxybenzoic acid)]

[] []

-Methylimidazole

[] []

Methylene blue

UiO--NH

Zr

-Aminoterephthalic acid (NH-HBDC)

Perylene

EuL

Eu+

[,′-(Carbonylbis (azanediyl))bis(methoxybenzoic acid)]

Pyrromethene / + sulforhodamine 

ZIF-

Zn+

-Methylimidazolate

[]

Rhodamine G

MIL-

Zn+

(NHMe)(tpt)(TZB)] (DMF)(tpt, ,,-tri(-pyridyl)-,,triazine; TZB, -(H-tetrazol--yl) benzoic acid)

[]

[]

Chapter 2 Guest-induced luminescence of metal-organic frameworks

27

Table 2.4 (continued) Organic phosphor

Host MOF

Rhodamine G

([Eu Eu+ (,-BDC).(phen) (HO)]n)

Rhodamine B

ZIF-

Rhodamine B Rhodamine B Rhodamine B

MOF- ZIF- Zr-MOF

Metal Linker centre

Zn+ +

References

,-BDC (BDC, terephthalic acid), ,-phenanthroline

[]

Imidazole--carboxyaldehyde

[]

Fumaric acid

[]

+

,-Dichloroimidazole (DCIM)

[]

+

,,-Benzenetribenzoic acid (BTB)

[]

Zr

Zn Zr

+

Rhodamine B

DUT-

Zr

Naphthalene-,-dicarboxylic acid

[]

Rhodamine B

ZIF-

Zn+

-Methylimidazole

[]

Such intense light can be obtained in two ways. The first consists of mixing red, green, and blue colours originating from separate sources (RGB method). The main drawback of this method is the difficulty in completely mixing these different colours. The second method of producing white light overcomes this drawback as it consists of one monochromatic source in which there is either the down-conversion of blue light or a combination of blue light and a yellow light – yellow light can be emitted by a phosphor. The latter mechanism can easily be achieved using composite materials consisting of yellow-emitting guests encapsulated in convenient blue-emitting MOFs to form a single component enclosing different phosphors. An example of white-light-induced emission within MOFs has been described by Mondal et al. [117] using CdTe quantum dots in [Cd (PyTa)(cba)2(MeOH)(H2O)] (PyTa, pyridylthioacetate; cba, carboxyvinyl benzoic acid). The encapsulation of CdTe was carried out using an ex situ method and the doped MOF exhibited a broad emission between 350–700 nm with two peaks observed at 412 (MOF assigned) and 563 nm (CdTe QD assigned) under excitation at 330 nm. A visible intense white light emission with CIE coordinates of (0.33, 0.32) was observed (Figure 2.13) [117]. Inspired by the RGB approach, Song et al. [105] designed a white LED (WLED), upon doping ternary Eu0.05Tb0.95BPT MOF (H3BPT, biphenyl-3,4′,5-tricarboxylate) with an organic dye. Upon the in situ encapsulation of blue-emitting coumarin 460 as a guest within the MOF comprised of Eu3+ (red emission) and Tb3+ (green emitting) ions, white emission was obtained with an overall quantum yield of 43.42%, which may be further applied in the lighting field. Depending on the external groups, size, and volume of their pores, MOFs are unique versatile materials that can accommodate several guest compounds, in addition to their special features. This functionality has been exploited by Tang et al. [109], who developed a promising approach to fabricate a highly stable WLED using a ZIF-8-like MOF, with Zn2+ as the metal centre and 2-methylimidazolate for the encapsulation of three organic dyes.

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

Figure 2.13: (a) Photoluminescence spectrum obtained for CdTe functionalised CP1 under excitation at 330 nm and a photographic image of the white light emission after illumination (inset). (b) The CIE diagram (republished with permission of Royal Society of Chemistry from reference [117]; permission conveyed through Copyright Clearance Center, Inc.).

Figure 2.14: A schematic illustration of the design and synthesis of polyurethane-coated MOF phosphor ZIF-8 pm546/pm605/SRh101 and the packaging of a remote-type WLED (republished with permission of Royal Society of Chemistry from reference [109]; permission conveyed through Copyright Clearance Center, Inc.).

The organic dyes, pyrromethene 546 (green emission), pyrromethene 605 (yellow emission), and sulforhodamine 101 (red emission), were encapsulated into the framework via an in situ self-assembly process at room temperature to obtain a blue-

Chapter 2 Guest-induced luminescence of metal-organic frameworks

29

Figure 2.15: A schematic representation of the encapsulation of [Ir(ppy)2(bpy)] in the MOF host, [(CH3)2NH2]15[(Cd2Cl)3(TATPT)4] · 12DMF · 18H2O, leading to white light emission (reprinted with permission from Springer Nature: reference [49]. Copyright 2013).

excitable broadband nanophosphor. To effectively prevent the phosphor from the disturbance of ambient oxygen and moisture, the as-synthesised phosphor was coated with optically transparent thermoplastic polyurethane (TPU) as a protecting and supporting matrix, and then moulded as a hemispherical phosphor layer in a remote configuration (Figure 2.14). A significant improvement in the fluorescence of 52.54% following the TPU coating was observed. With the high plasticity of TPU, a stable remote-type white-lighting device was fabricated with a blue LED chip, which exhibited similar-to-incandescent performance with CIE coordinates of (0.465, 0.413) and a low CCT of 2,642 K. This indicated the considerable potential of the phosphor-MOF composite for highly stable warm indoor illumination, and that it was a promising candidate for replacing remaining incandescent bulbs [109]. Another interesting example of a white-light emitting composite was prepared by Sun et al. [49] using the combination of a blue-light-emitting MOF, encapsulating an Ir3+ complex, which emits yellow light. The construction of the blue-emitting MOF was achieved using a hexadentate carboxylate triazine ligand (2,4,6-tris(2,5-dicarboxylphenylamino)-1,3,5triazine), and cadmium ions as the metal centres. The resulting MOF, [(CH3)2NH2]15[(Cd2Cl)3 (TATPT)4] · 12DMF · 18H2O, emits light at 425 nm, which was attributed to the fluorescent organic TATPT linker. [Ir(ppy)2(bpy)]+ (ppy, 2-phenylpyridine; bpy, 2,2′-bipyridne) was

30

Moritz Bassen, Katia Nchimi, Huayna Terraschke

then loaded into the pores of the MOF, leading to the emission of bright white light when excited at 370 nm with a quantum yield of 20.4% (Figure 2.15) [49]. Aiming at the single-phase production of WLEDs, down-conversion can also be supported by the guest@MOF luminescence. An efficient down-converter was designed by Gutiérrez et al. [48] by coating Gaq3@ZIF-8 MOF onto the surface of a 405 nm LED. The LED lamp was dipped in a dispersion of the Gaq3 complex ([gallium (III) tris(8-hydroxyquinolinato)]) in a 10% w/w solution of transparent thermoplastic polymer (PMMA) in dichloromethane. The ZIF-8 MOF was loaded with the gallium complex, and the subsequent white luminescence originated from the combination of the violet light (405 nm) of the LED and the green-yellow emission of the coating (Figure 2.16). Moreover, the encapsulation of the dye within the MOF enhanced the longterm stability of the nanomaterial, which was found to be stable for at least 8 months.

2.4.2 Chemical sensing of ions Chemical sensing consists of the selective and sensitive detection of chemical species (ions, complexes, or organic molecules), upon the recording of an analytically useful signal. The appearance, disappearance, or change in the luminescence properties of MOFs induced by the investigated species in the guest–host interaction is a reliable tool for assessing and monitoring the presence of analytes. Among species for which guest@MOFs exhibit significant real potential application in sensing, anions (such as F–) or cations (such as Fe3+ or Cu2+) occupy a great place.

Figure 2.16: (A) Emission spectrum obtained for the 405 nm LED coated with Gaq3@ZIF-8 (0.5 mmol) dispersed in a PMMA polymer matrix. The inset is a representation of the CIE (1936) coordinates of the device (0.27, 0.34). (B) Photograph of the turn-off (left) and turn-on (right) down-converter MOF-LED device (reproduced with permission from reference [48]).

Chapter 2 Guest-induced luminescence of metal-organic frameworks

31

Chemical sensing of fluoride ions Fluoride is an essential trace element for the human body, but it can cause serious health problems, including dental and skeletal fluorosis, when overdosed, and it is one of the most serious pollutants in water because of its high toxicity. This makes sensing techniques very important. Zheng et al. [40] reported that Tb3+-doped MOFs can be used to sense fluoride ions. Zr-MOF was constructed using Zr4+ metal centres, terephthalic acid, and isophthalic acid as linkers. Using post-synthetic modification, Tb3+ ions were encapsulated inside the cavities of the host MOF. A strong and broad emission band was observed, with peaks at 488, 544, 584, and 620 nm. The composite materials, loaded with Tb3+ ions, were dispersed in aqueous solutions containing aliquots of different anions, including Cl−, Br−,I−, S2−, NO3−, NO2−, CO32−, HCO3−, SiO32−, SO42−, and PO43−. In the presence of F–, a drastic increase in the MOF emission was observed, while the other tested anions did not show any significant changes. Such modifications can even be monitored using the naked eye, underlining the good selectivity towards the detection of fluoride ions (Figure 2.17). The authors have suggested that this selectivity towards fluoride anions can be explained by the strong nucleophilic ability of F− and its binding to Zr4+sites within the MOF [40].

Chemical sensing of iron(III) ions Iron is one of the most abundant elements found on the earth, and in the form of metals, ions, or oxides, a large amount of its compounds is used in various fields ranging from metallurgy to medicine. In the human body, iron is mostly found in the form of Fe(III) ions, and it is one of the trace elements that performs vital functions. However, Fe concentrations should remain within specific thresholds because Fe(III) deficiency or excess can be harmful and may lead to pathologies, such as cell death or Parkinson’s disease [118]. Iron-sensing can be supported by MOFs, as described in the following two examples. One of the examples we have selected consists of a lanthanide-doped MIL-53-COOH (Al) MOF. Eu3+ ions were encapsulated inside the pores of the MIL-53-COOH(Al) nanocrystals, affording an optically active composite. The MOF was solvothermally synthesised from a mixture of aluminium chloride, trimellitic acid, and DMF. The Eu3+ ions were then encapsulated into the MOF by soaking the as-prepared MOF in an ethanolic solution of the Eu3+ chlorate salt for two days. Later, the MOFs were washed to remove any residual Eu3+ ions from their surface. Upon mixing Eu3+@ MIL-53-COOH(Al) with increasing amounts of Fe3+, an intense quenching of the emission peaks centred on Eu3+ (Figure 2.18) was observed, which is most likely due to cation exchange inside the framework of MIL53-COOH (Al) to form MIL-53-COOH (Fe). The emission of the Eu3+-doped Al-based MOF in the presence of Fe3+ ions was compared with that of the same material in the presence of competing metal ions, including abundant mineral salts (K+, Na+, Mg2+, Ca2+, and Cu2+),

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

Figure 2.17: Emission spectra obtained for Tb3+@Zr-MOF dispersed in different anion solutions (excited at 302 nm). The inset is a photograph of the corresponding samples under UV-light irradiation (reprinted with permission of American Chemical Society from reference [40]).1

oligo-elements (Zn2+, Co2+, Ni2+, Fe2+, and Ag+), and toxic metal ions (Hg2+, Cd2+, Pb2+, and Cr3+). Among all the ions tested, Fe3+ showed the most significant quenching effect on luminescence. The quenching was attributed to cation exchange of the Al3+ cation as the central metal ion in the MOF with Fe3+ ions, which was able to transform MIL-53-COOH (Al) into MIL-53–COOH (Fe). Within the newly formed MOF, a ligand-to-metal charge transfer was observed by the suppression of the sensitisation of the Eu3+ cations via the antenna effect. Good selectivity towards morphological and optical features makes such an Eu3+-doped MOF applicable for the sensing of Fe3+ ions [31] The second example is based on the work of Xu et al. [32], who showed a comparable approach using Eu3+-doped MIL-24 MOF to sense different types of metal ions. In addition to the quenching of the MOF luminescence due to Fe3+ ions, which results in an extremely low emission intensity, Fe2+ ions were also detected. While the Fe3+ ions quenched the emission of the whole MOF by replacing the Ga3+ ions in the framework, the Fe2+ ions quenched the red emission of the Eu3+ ions by substituting them as host molecules within the MOF pores (Figure 2.18). Therefore, the emission colour of the MOF was blue-shifted. In addition, the intensity is significantly lower than that without the presence of Fe2+ ions but is still visible (Figure 2.19) [32].

 https://pubs.acs.org/doi/10.1021/acsomega.8b02134 further permission related to the material excerpted should be directed to the ACS.

Chapter 2 Guest-induced luminescence of metal-organic frameworks

Figure 2.18: Response of the fluorescence of Eu3+@MIL-53-COOH (Al) towards aqueous solutions of various metal cations (republished with permission of Royal Society of Chemistry from reference [31]; permission conveyed through Copyright Clearance Center, Inc.).

Figure 2.19: Luminescence intensity of Eu3+@MIL-124 when interacting with different metal ions in an aqueous solution (reprinted with permission from reference [32]. Copyright 2015 American Chemical Society).

33

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Moritz Bassen, Katia Nchimi, Huayna Terraschke

Chemical sensing of copper(II) ions Lin et al. [60] have reported a highly luminescent CQD@ZIF-8 composite that is optically responsive in the presence of Cu2+ ions. For the preparation of the material, branched polyethyleneimine (BPEI) was used as a capping agent during the one-step synthesis of the CQDs to afford the BPEI-CQD@ZIF-8 composite, in order to ensure good dispersion and an even distribution of the nanoparticles. Upon excitation at 365 nm, the obtained material showed a strong blue emission. However, upon the addition and adsorption of Cu2+cations into the ZIF-8 matrix, significant quenching of the emission of the CQD-loaded MOF occurred (Figure 2.20) [60]. This quenching effect was higher for the composite than for BPEI-CQDs alone, indicating the synergistic effect of accumulation and quenching. Moreover, the composite shows high selectivity towards Cu2+ with no further effect on the quenching behaviour in the presence of the other metal cations, anions, or organic components studied. The positive selectivity towards Cu2+ over the anions may be attributed to the electrostatic repulsion between the anions and the anionic ZIF-8 MOF in the composite. Moreover, the authors suggested that the synergetic effect of the selective adsorption of heavy metal ions by ZIF-8 MOF and the selective sensing of Cu2+ by the BPEI-CQDs makes the composite MOF very useful for sensing Cu2+ ions [60].

Figure 2.20: Emission spectra (a, b), excitation spectra (c, d), and UV absorption spectra (e, f) of the BPEICQD@ZIF-8 solutions in the absence (a, c, e) and presence (b, d, f) of Cu2+ ions. The inset shows the dispersions of BPEI-CQD@ZIF-8 in the absence (left) and presence (right) of Cu2+ ions (reprinted with permission from reference [60]. Copyright 2014 American Chemical Society).

Another interesting material for Cu2+ ion sensing is lanthanide-doped MOF Zn-BETC reported by Luo et al. [25], and previously described in Section 3.1. When the MOF was doped with Eu3+, the authors observed that the infusion of the composite with variable amounts of metal salts resulted in significant differences in the resulting emission spectra. Indeed, Na+, K+, and Zn2+ ions have a negligible effect on (Figure 2.21) the

Chapter 2 Guest-induced luminescence of metal-organic frameworks

35

luminescence intensity, whereas Cu2+, Co2+, Mn2+, and Ni2+ showed different ranges of quenching of the luminescence intensity of the Eu3+-loaded MOF. The quenching effect was attributed to the replacement of NH4+ counterions within the channels of the MOF by transition metal ions. More specifically, the emission centred on Eu3+ in the presence of a 10–6 M solution of CuCl2 was 16 times higher than that in the presence of a 10–2 M solution of the same copper(II) chloride salt. This makes the described system potentially interesting for low-detection applications [25].

Chemical detection of nicotine Yan et al. [108] designed a MOF composite loaded with methylene blue (MB) for sensing nicotine in urine and in living cells. This is interesting as nicotine is genotoxic and tumour-promoting and is excreted via urine. The host MOF, UiO-66-NH2, was synthesised using a hydrothermal reaction of ZrCl4 and 2-aminoterephthalic acid (NH2-H2BDC). The cyan colour complex was prepared by loading MB into the MOF using a conventional impregnation method. MB@UiO-66-NH2 does not show any emission at ~ 700 nm in aqueous systems, which is the typical emission of free MB. Upon excitation at 357 nm, MB@UiO-66-NH2 shows a distinct emission at 449 nm, which was assigned to the ligand to metal charge the transfer process. As urine is an aqueous solution containing a variety of components, mainly creatinine, urea, SO42–, Na+, K+, NH4+, and glucose, the optical response of the MB@UiO-66-NH2 with nicotine was compared to those of the above-mentioned analytes. Only in the presence of nicotine did the luminescence intensity of MB@UiO-66-NH2 exhibit a >20-fold enhancement in magnitude, while the intensity was almost unchanged in the presence of the other analytes studied. This discrepancy can be visualised using the naked eye, which paves a direct and simple method for unique nicotine recognition. The mechanism of this selective recognition of nicotine involves the transformation of MB into leucomethylene blue (LB) in the presence of a reducing agent. Nicotine acts as a reducing agent, and serves as a hydrogen acceptor. The LB formed can be reoxidised by oxygen in the air, in terms of light excitation, and serves continuously as a hydrogen acceptor. The remarkable enhancement in the intensity and obvious blue shift in the emission wavelength were attributed to the photoinduced electron transfer from LB to MOF UiO-66-NH2 [108]. Numerous other guest-encapsulated MOF composites have been synthesised for sensing ions, organic molecules, and water by many, including Li et al. [119] (ClO– and SCN–), He et al. [66] (CO2), Gao et al. [82] (H2S), Takashima et al. [120] (aromatic VOCs), Yang et al. [114] (Fe3+, Cr3+, Ga3+, pesticides, and nitro explosives), Zhang et al. [115] (Fe3+ and Cr2O72–), and Zhang et al. [121] (picric acid).

36

Moritz Bassen, Katia Nchimi, Huayna Terraschke

Figure 2.21: A schematic representation of the proposed mechanism of nicotine sensing in MB@UiO-66NH2 (reprinted with permission from reference [108]. Copyright 2019 American Chemical Society).

2.4.3 Other applications Photocatalytic oxidation The oxidation of alcohols into carbonyl compounds (aldehydes, ketones, or carboxylic acids) is an important reaction in organic synthesis. To mitigate the environmental cost of such transformations, there is a revived interest in the development of selective and effective catalysts. A luminescent composite for this photocatalytic organic transformation was engineered by Ke et al. [122], which was used for the selective oxidation of aromatic alcohols. To this end, an MIL-100 (Fe)-like MOF was synthesised from FeCl3 ·6H2O and benzene-1,3,5-tricarboxylic acid (H3BTC). Later, the MOF was dispersed in DMSO, and Cd(CH3COO)2 · H2O was added to form CdS quantum dots in a Teflon-lined autoclave. The resulting CdS@MIL-100 (Fe) composite exhibited excellent photostability and enhanced photocatalytic activity with respect to the selective oxidation of benzyl alcohol to benzaldehyde under visible light irradiation. A series of samples were synthesised by adding different amounts of MIL-100 (Fe) to the toluene reaction solution, which resulted in a different but overall significantly higher conversion rate, under excitation at 420 nm (Figure 2.22) [122].

Phototherapy In another interesting example, Zhou et al. [67] focused on the synthesis of porphyrinic MOFs, encapsulating gold nanorods (AuNRs), used for the combinational phototherapy of cancer. The potential of core–shell AuNR@MOF heterostructures containing individual AuNRs as the core and mesoporous porphyrinic MOFs as the shell in combinational

Chapter 2 Guest-induced luminescence of metal-organic frameworks

37

Conversion (%)

50 40 30 20 10 0 (a)

(b)

(c)

(d)

(e)

(f)

(g)

(h)

Photocatalyst Figure 2.22: Conversion rates for the photocatalytic selective oxidation of benzyl alcohol to benzaldehyde in the presence of different photocatalysts under visible light irradiation (λ > 420 nm) for 5 h: (a) MIL-100 (Fe), (b) CdS, (c–h) increasing concentration of CdS@MIL-100 (Fe) (republished with permission of the Royal Society of Chemistry from reference [122]; permission conveyed through Copyright Clearance Center, Inc.).

photodynamic therapy (PDT) and photothermal therapy (PTT) against tumours was examined (Figure 2.23). Because of their outstanding photothermal conversion ability, AuNRs have been widely used as photothermal agents. Porphyrinic MOFs, such as PCN222, were used to demonstrate that combining these MOFs with plasmonic AuNRs in a single heterostructure enables efficient wavelength-dependent light harvesting for photothermal conversion, and singlet oxygen production for tumour therapy. AuNR@MOF showed greater antitumor efficiency than PDT or PTT alone [67].

2.5 Concluding remarks This chapter provides an overview of the guest species encapsulated in the pores of MOFs to enhance their luminescent properties. Using lanthanide ions, coordination complexes, quantum dots, plasmonic nanoparticles, or organic phosphors, a vast range of guest-doped MOF composites have been described and examined, which show significant effects on their optical properties. For example, Zn-based bio-MOF1 was functionalised into Ln3+@bio-MOF1, a material exhibiting new and tuneable emission features, upon doping with various lanthanide ions, such as Sm3+, Eu3+, Tb3+, and Yb3+ [30]. In other cases, the starting MOF is already emissive, but the introduction of guest molecules enhances the luminescent properties, resulting in new functionalities within the MOF. For instance, Sun et al. developed a white light-emitting material upon combination of a blue-emitting MOF with a yellow-emitting Ir3+ complex [49]. Indeed, it has been shown that these materials are especially efficient for lighting, but they can also

38

Moritz Bassen, Katia Nchimi, Huayna Terraschke

Figure 2.23: (a) A schematic illustration of the synthesis of AuNR@MOFs and (b) use of AuNR@MOFs for dual-light-activated combinational phototherapy (republished with permission of Royal Society of Chemistry from reference [67]; permission conveyed through Copyright Clearance Center, Inc.).

be used for sensing ions or small organic molecules via quenching or shift of emission through different mechanisms. A Eu3+-doped MIL-53_COOH(Al) MOF was prepared by Zhou et al. [31], which showed remarkable selectivity for the detection of Fe3+ ions via the quenching of the luminescence of Eu3+. Lan et al. reported a CQD@ZIF-8 MOF with blue luminescence, which was only quenched in the presence of Cu2+ ions [55]. Fluoride anions are easily detected even with the naked eye using a Tb3+@Zr-MOF composite through the tremendous increase in the Tb3+ emission intensity [40]. On the other hand, the guest-induced luminescence of MOFs can also be utilised for photocatalytic applications. The functionalisation of MIL-100 (Fe)-like MOF with CdS quantum dots developed by Ke et al. [122] showed improved conversion rates for the oxidation of benzyl alcohol to benzaldehyde under visible light irradiation. Although the encapsulation of guest species into MOFs overcomes some of the issues observed in terms of the luminescence of single luminescence species, such as quenching via the aggregation of the materials, the encapsulation process still entails some important problems that need to be overcome. The main one is to ensure an even distribution of the guest species over the entire MOF, which will significantly improve

Chapter 2 Guest-induced luminescence of metal-organic frameworks

39

the efficiency of the MOF luminescence. In general, it should be the aim to achieve improved efficiency to make these materials attractive for large-scale applications, such as the lighting of residences or to fulfil the high standards of display technology. Many highly promising approaches have already been developed, but progress is still required.

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Patric Lindenberg, Katia Nchimi, Huayna Terraschke✶

Chapter 3 Synthesis, functionalisation, and applications of iron oxide magnetic nanoparticles Magnetic nanomaterials are important for potential applications in several fields, such as nanocatalysis, drug delivery, biosensing, and magnetic resonance imaging contrast agents. This chapter not only offers an overview of the most important types of magnetic nanoparticles and how they are typically synthesised and functionalised, but also discusses the problems and the future of this exciting research field.

3.1 Introduction The most applied materials for magnetic nanoparticles (MNPs) synthesis are iron oxides, which can exist as hematite (α-Fe2O3), maghemite (λ-Fe2O3), and magnetite (Fe3O4). Blood-red hematite (Greek haima represents blood) is the oldest known iron oxide mineral, widely found in nature in rocks and soils, and an important pigment and valuable ore [1–3]. In nature, red-brown maghemite (magnetite + hematite) is found in soils, typically as a product of heating other iron oxides or as a weathering product of magnetite

Figure 3.1: Natural magnetite crystal extracted in the Goiás State in Brazil (reproduced with permission from Marvin Radke).

✶

Corresponding author: Huayna Terraschke, E-Mail: [email protected]

https://doi.org/10.1515/9783110459098-003

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[3]. Similarly, black magnetite (Figure 3.1) is also an important iron ore that is, in combination with titanomagnetite, responsible for the magnetic properties of rocks [3]. There are two main groups of chemical routes for the preparation of nanomaterials, in general, and also of iron oxide MNPs, namely, bottom-up and top-down approaches [1]. Bottom-up processes are the most commonly applied, and include co-precipitation, thermal decomposition, hydrothermal and solvothermal synthesis, sol–gel synthesis, the polyol method, microemulsion, sonochemical synthesis, and microwave-assisted synthesis [4, 5]. These approaches are advantageous because of the well-controlled particle size and shape of the products, which is enabled by flexible changes in the experimental parameters, such as temperature or concentration, often resulting in the production of monodispersed nanoparticles [1]. Top-down methods, including various types of lithography, milling, or etching, are less common and are therefore not contemplated in detail within this chapter [1]. In general, MNPs are important for applications in diverse medical and technological fields such as nanocatalysis, biosensing, magnetic storage, electrochemical devices, and electronic memory devices [1, 2, 6, 7]. In medicine, the use of MNPs as contrast agents for magnetic resonance imaging (MRI) is particularly interesting, as is their use in new hyperthermia therapy, which use denaturation of cancer cells by heat treatment. Furthermore, MNPs can be used for drug delivery, owing to their thermosensitivity [8]. For medical and biological applications, the functionalisation of the magnetic nanomaterial surface plays a crucial role because of the hydrophobic properties of most MNPs, and the consequent disadvantages for their use in aqueous environments, such as in the human body. Moreover, the stability of these particles are of great interest for these types of applications. Therefore, these characteristics should be known and controlled, which can be achieved by coating iron nanoparticles with different functionalising shells [7, 9]. This chapter offers a summary of the vast literature on the most applied materials for the production of MNPs: iron oxides. In detail, Section 3.2 describes the types and crystal structures of these raw materials, while Section 3.3 explains the most common synthesis methods for MNPs, as well as their disadvantages, their advantages, and strategies for controlling their particle size and morphology. Section 3.4 highlights recent studies on the formation mechanism of iron oxides. Section 3.5 focuses on important applications of MNPs, which are achieved by the advanced surface functionalisation techniques explained in Section 3.6. For further reading, the books Magnetic Nanoparticles: From Fabrication to Clinical Applications [10], Magnetic Nanoparticles: Particle Science, Imaging Technology, and Clinical Applications [11], and Magnetic Nanoparticles [12] are highly recommended.

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3.2 Types of iron oxide magnetic nanoparticles Among the nine types of iron oxides, the most applied for the production of MNPs are hematite (α-Fe2O3), maghemite (λ-Fe2O3), and magnetite (Fe3O4) (Figure 3.1) [13]; hematite and maghemite are iron(III) oxides, and magnetite is a mixed valence iron(II,III) oxide compound. Hematite is the most stable iron oxide, and it is broadly applied in catalysts, pigments, and gas sensors, particularly because of its low cost. This compound exhibits weak ferromagnetism at room temperature [13], and is an n-type semiconductor with a bandgap of 2.3 eV [13]. The crystal structure of α-Fe2O3 was originally reported by Bragg and Bragg [14] and Pauling and Hendricks [15] at the beginning of the twentieth century. It is isostructural with corundum, comprising a hexagonal unit cell (a = 0.5034 nm, c = 1.375 nm) with hexagonal close-packed arrays of oxygen ions along the [001] direction, in which the anion planes are parallel to the (001) plane. Two-third of the sites are filled with Fe3+ ions, presenting a regular arrangement in the (001) plane, in which two filled sites are followed by one vacant site, resulting in the formation of sixfold rings. In addition, pairs of FeO6 octahedra are formed. Each octahedron shares edges with three other octahedrons and one face with a neighbouring octahedron in an adjacent plane. The face-sharing feature occurs along the c-axis, resulting in the distortion of the cation sublattice from ideal packing. Within the hematite structure, the oxygen–oxygen distances along the shared face are shorter than those in the unshared faces and are therefore trigonal distorted. The resulting Fe–O–Fe triplet arrangement around the shared face influences the magnetic properties of α-Fe2O3 [3]. Maghemite exhibits ferrimagnetism at room temperature [13]. Its crystal structure is very similar to that of magnetite, resulting in similar X-ray diffraction (XRD) patterns (Figure 3.2). The difference between these structures is that, in maghemite, the iron ions are in a trivalent oxidation state. Therefore, the cation vacancies in maghemite compensate for the divalent iron ions present in the magnetite structure. λ-Fe2O3 has a cubic unit cell, with a = 0.834 nm, where every unit cell comprises 32 O2– anions, 21⅓ Fe3+ cations, and 2⅓ vacancies. In this unit cell, eight cations occupy tetrahedral sites, while the other cations are randomly dispersed over the octahedral sites, and the vacancies occupy the octahedral sites [3]. Figure 3.2 shows the simulated XRD patterns of hematite [14], maghemite [16], and magnetite [17]. In comparison to magnetite, the Bragg reflections of maghemite are slightly shifted to higher angles, and additional reflections are observed, for example, between 2θ of 10° and 30°. Similar to hematite, the structure of magnetite was investigated by Bragg et al. [17] at the beginning of the twentieth century. Magnetite was one of the first mineral structures investigated by XRD and was originally reported in 1915 [3, 17]. This structure consists of an inverse spinel (Figure 3.3), with a face-centred cubic unit cell (a = 0.839 nm) containing 32 O2– anions, which are regularly cubic close packed along the [111] direction, with eight formula units per unit cell. The general frequently non-stoichiometric formula of magnetite is Y[XY]O4 (X = Fe2+, Y = Fe3+), in which the brackets represent the octahedral sites. In

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Figure 3.2: XRD patterns of hematite (α-Fe2O3) [14], maghemite (λ-Fe2O3) [16], and magnetite (Fe3O4) [17].

Figure 3.3: Structure of Fe3O4, [2, 17] important material for the fabrication of magnetic nanoparticles (reproduced with permission from reference [13]).

total, Fe2+ and Fe3+ ions are distributed over eight tetrahedral sites, and octahedral and mixed tetrahedral/octahedral layers are formed along the (111) direction [3]. As described by Shrivastava et al. [18], when the crystal size is reduced to the nanoscale, magnetite presents a long-range ordering of the magnetic moment, high surface-tovolume ratio, and lower toxicity in comparison to the bulk analogue material [18]. In addition, iron oxide nanoparticles are superparamagnetic at room temperature for dimensions

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below 15 nm, at which the thermal energy can overcome the anisotropic energy barrier for single particles. In particular, one-dimensional iron oxide nanostructures exhibit unique physicochemical properties, owing to their high intrinsic anisotropy and surface activity [13]. However, the application of special synthesis and functionalisation techniques is necessary to produce magnetic nanomaterials, with well-defined particle sizes and morphologies, to control their magnetic properties [19].

3.3 Synthesis methods Among the most important issues to be considered during the synthesis of MNPs are achieving a monodispersed particle size distribution, and controlling the shape of the nanostructure. In addition, it is desirable to achieve a satisfactory degree of crystallinity, naturally, of the desired modification. Finally, the produced magnetic nanomaterials must be stable over a long period of time [1]. To achieve these objectives, the most commonly applied synthesis methods for iron oxide nanomaterials are co-precipitation, thermal decomposition, hydrothermal and solvothermal synthesis, sol–gel synthesis, the polyol method, microemulsion, sonochemical synthesis, microwave-assisted synthesis, and biosynthesis [13], which are briefly described below. Co-precipitation is the simplest synthesis method. For example, for the preparation of magnetite nanoparticles, the typical approach involves a mixture of stoichiometric amounts of divalent and trivalent iron salts under an inert atmosphere in a very basic solution. In this case, pH values below 11 favour nucleation, while pH values above 11 favour crystal growth [13]. An additional relevant consideration is the application of mechanical and non-magnetic agitation to avoid aggregation of the produced nanoparticles around the magnetic stirrer. This approach can yield gram-scale production, in which the simplified reaction mechanism can be summarised as [13, 20] Fe2+ + 2Fe3++ 8OH− ⇄ FeðOHÞ2 + 2FeðOHÞ3 ! Fe3 O4 # + 4H2 O

(3:1)

There are different adaptations of co-precipitation synthesis, such as applying ultrasound to obtain ca. 15 nm particles and applying surfactants and biomolecules to improve biocompatibility and avoid aggregation [13]. Examples of typically employed organic compounds and biomolecules include octanoic acid and chitosan, which improve the dispersibility in liquid media. Coating Fe3O4 nanoparticles with chitosan can be achieved by substituting water during synthesis with a 2% chitosan solution in acetic acid [13]. Despite the simplicity and mild conditions required by the co-precipitation approach, this strategy also presents a few disadvantages, such as fast crystallisation and limited control over particle size and morphology, generally yielding a broad particle size distribution. These factors depend on experimental parameters such as the applied iron salts, Fe2+/Fe3+ ratio, pH value, and ionic strength of the solvent [13, 20].

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The thermal decomposition of precursors is one of the most applied non-aqueous strategies for the synthesis of magnetic nanomaterials. This method can be divided into two main approaches. In the first approach, the precursors are added to a hot reaction mixture (hot injection), while in the second approach, the precursors are first prepared at room temperature and then heated posteriorly [13]. The higher crystallinity and narrower size distribution of the resulting nanomaterials are the most important advantages of this strategy. A typical synthesis involves the decomposition of coordinated iron precursors in organic solvents. Examples of frequently used iron precursors include [Fe(CO)5] (Figure 3.4), [Fe(acac)3] (acac, acetylacetone), iron oleate, [Fe(Cup)3] (Cup, N-nitrosophenylhydroxylamine), Prussian blue, iron urea complexes, ferrocene, and [Fe3(CO)12]. In addition, oleic acid, 1-octadecene, 1-tetradecene, and oleylamine are examples of organic molecules frequently used as stabilisers, which favour a slow nucleation process and hinder the adsorption of additives on the crystal surface, thus

Figure 3.4: Schematic representation of the synthesis of iron oxide nanoparticles by thermal decomposition (reproduced with permission from reference [21]).

Figure 3.5: Schematics of hydrothermal and solvothermal procedures for the preparation of α-Fe2O3 and Fe3O4 nanoparticles, respectively (reprinted from reference [22], with permission from Elsevier).

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inhibiting the growth of nanoparticles. In this method, the shape and size of the nanoparticles are tuned by the temperature, precursors, additives, and solvents [13]. Hydrothermal and solvothermal syntheses are wet-chemical synthesis techniques in which aqueous and non-aqueous solutions, respectively, are heated above their boiling point under a high vapour pressure in a sealed container [13]. For hydrothermal synthesis, water is the applied reaction medium, while various organic solvents are used as the reaction medium for the solvothermal approach [22]. In the classic solvothermal reaction for the production of iron oxide nanoparticles, Fe3+ is applied as an iron source, and acetate, urea, and sodium citrate are mixed in ethylene glycol and sealed in a Teflon-lined autoclave at ca. 200 °C for 8–24 h. Hydrothermal and solvothermal strategies are widely applied to obtain highly crystalline hematite, maghemite, and magnetite nanoparticles [13]. Using these strategies, the nature, size, morphology, and crystallisation behaviour of the final products can be controlled by varying the polarity, surface tension, viscosity, and coordination ability of the solvents [22]. For example, Su et al. [22] prepared α-Fe2O3 and Fe3O4 nanocrystals at 200 °C under pressure (Fig. 3.5). α-Fe2O3 was obtained via hydrothermal synthesis, and Fe3O4 was obtained under solvothermal conditions in ethylene glycol, where urea was added to favour precursor formation [22]. The urea concentration was varied to tune the size and morphology of the obtained particles. Hence, the shape of the α-Fe2O3 particles changed from olive-like to rhomb-like structures, and the shape of the Fe3O4 particles varied from hollow spheres and pinecone-like structures to cracked nanostructures [22]. Many iron oxide nanoparticles have also been synthesised by the sol–gel approach, in which “sol” refers to the formed stable dispersion of colloidal particles or polymers in a solvent, and “gel” refers to the formed three-dimensional continuous network enclosing the liquid phase, which is generally divided into colloidal or polymer gels [13]. Within colloidal gels, the network is formed by the aggregation of nanoparticles interacting through van der Waals forces and hydrogen bonds. In polymer gels, the obtained particles present a polymeric substructure formed by the aggregation of subcolloidal nanoparticles. In this case, the gel is formed by linking of the polymer chains. Examples of typically applied precursors include iron alkoxides and salts, such as chlorides, nitrates, and acetates, which undergo hydrolysis and polycondensation reactions. This procedure is generally performed at room temperature, with posterior heating, to obtain the final product [13]. In the so-called polyol method for the production of iron oxide particles, polyols are used simultaneously as solvents, reducing agents, and stabilisers to control particle growth and avoid particle aggregation [13]. In a typical synthesis procedure, an iron precursor, such as [Fe(acac)3], is suspended in a liquid polyol, where ethylene, diethylene, triethylene, or tetraethylene glycol are the most commonly used. Subsequently, the mixture is stirred and heated close to the polyol boiling point. The advantages of this strategy include the following: (i) a high pressure is not required, (ii) the particle surface is coated in situ with hydrophilic ligands and can, therefore, be easily suspended in aqueous solutions or other polar solvents, and (iii) the high reaction temperature favours a high crystallinity. One disadvantage of the polyol method is its high energy and material costs [13].

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Figure 3.6: Field-emission scanning electron microscopic (FE-SEM) images of a grass-like Fe3O4 nanostructure produced by combined microemulsion and solvothermal synthesis (reprinted from reference [23], with permission from Elsevier).

Microemulsions are liquid mixtures of oil, water, and surfactants. The surfactants are responsible for forming a monolayer at the interface between oil and water, in which the hydrophobic tail is dissolved in the oil phase and the hydrophilic head groups are dissolved in the aqueous phase [13, 24]. In general, the aqueous phase contains metal salts and other reagents, while the oil phase consists of a complex mixture of hydrocarbons and olefins [13, 24]. There are different types of microemulsions, for example, direct and reverse microemulsions. In direct microemulsions, the oil is dispersed in the aqueous phase (o/w), while in reverse microemulsions, water is dispersed in the oil phase (w/o). The surfactants typically used in these methods are bis (2-ethylhexyl) sulfosuccinate (AOT), sodium dodecyl sulphate (SDS), cetyltrimethylammonium bromide (CTAB), and polyvinylpyrrolidone (PVP). The particle size can be controlled by varying the droplet size, reactant concentration, and surfactant type [13]. An additional advantage of this method is its good performance at room temperature. However, the major disadvantages are the several washing steps required to remove the surfactants and the strong agglomeration of the produced nanoparticles [13, 24]. Li et al. [23] applied a combination of a reverse microemulsion and solvothermal approach to produce grass-like Fe3O4 nanostructures (Figure 3.6). The reverse microemulsions were prepared by mixing a solution of FeCl3 in ethylene glycol and a microemulsion composed of a KOH aqueous solution in a mixture containing the surfactant Triton™ X-100, n-octanol, and cyclohexane. The resulting reverse microemulsion was then transferred to a Teflon-lined stainless-steel autoclave and heated at 190 °C for 10 h. The grass-like Fe3O4 nanostructures were finally obtained after centrifuging, washing, and drying [23]. Sonochemical methods are based on acoustic cavitation, caused by high-intensity ultrasound radiation, in which alternating expansive and compressive acoustic waves cause the formation of oscillating bubbles (cavities) in a solution; they are often applied for the production of Fe3O4 nanoparticles. These bubbles accumulate ultrasonic energy, grow to a certain size, and collapse, releasing the accumulated energy within a very short period of time, and reaching temperatures of 5,000 K and pressures of up to 1,000 bar [13, 26]. Some advantages of this method include uniform mixing, synthesis under ambient conditions, and reduced crystal growth. The major disadvantage, however, is poor control

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Figure 3.7: Schematic of the growth mechanism of porous α-Fe2O3 nanostructures, prepared by a combination of solvothermal and microwave-assisted synthesis (reproduced from reference [25] with permission from the Royal Society of Chemistry).

of the shape and dispersity [13, 26]. For microwave-assisted strategies, molecules align their dipoles according to the external field, upon excitation with microwave radiation, resulting in strong agitation due to the reorientation of the molecules in phase with the electrical field excitation [13]. The advantages of this method include intense and homogeneous internal heating, reduced synthesis time, and consequent low energy cost. This method has been applied for the synthesis of magnetite, maghemite, and hematite nanoparticles, with reaction times of as low as 25 min [13]. One of the many applications of this method was reported by Rao et al. [25], who used a combination of solvothermal and microwaveassisted synthesis (Figure 3.7) to prepare porous α-Fe2O3 nanostructures for application as a negative electrode material in lithium and sodium batteries. In addition to these conventional methods for the synthesis of iron oxide magnetic nanocrystals, alternative methods, such as biomineralisation of these nanoparticles in living microorganisms (for example, bacteria), have been reported in the literature [13, 27, 28]. This method is a bottom-up approach, based on reduction and oxidation reactions, which explores the antioxidant or reducing properties of microbial enzymes or plant phytochemicals. Typical experiments are conducted with magnetotactic or iron-reducing bacteria, such as Actinobacter sp. or Bacillus sp., under aerobic or anaerobic conditions. The low environmental impact of the chemicals and process as well as the high biocompatibility of the produced MNPs are the major advantages of this strategy [13, 28]. The large variety of synthesis methods discussed above offers enhanced freedom in the morphology of the prepared nanocrystals. As summarised in Table 3.1, iron oxides (for example, Fe3O4) can be obtained not only as monodispersed spherical particles but also as octahedral particles, cubic particles, nanorods, nanowires, nanorings, and nanoprisms, to name only a few. In this context, surfactants and templates play important roles, as also temperature, ageing time, and reactant concentrations during solvothermal

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Table 3.1: Examples of synthesis strategies for Fe3O4 magnetic nanocrystals with different geometries [2]. Geometry

Synthesis method

Examples

Spherical

Polyol approach

Oleylamine applied as a reducing agent and stabiliser; tunable monodispersed particle size achieved by varying the ratio between benzyl ether and oleylamine []

Thermal decomposition

Iron(III) oleate applied as a precursor, thermally decomposed in liquid phase at  °C for  min; crystal size tuned between  and  nm by changing the decomposition temperature and ageing time []

Octahedral

Solvothermal synthesis

Potassium ferrocene, sodium borohydride and hydroxide, polyvinylpyrrolidone, and alcohol heated at  °C in an autoclave; growth of the () facet is minimised, owing to the adsorption of hydroxyl on the () facet, influencing the growth direction []

Cubic

Polyol approach

Comparison between different stabilisers (sodium oleate, potassium oleate, dibutylammonium oleate, and oleic acid) [] indicates that the growth rate anisotropy is related to the adhesion of the stabilising agent on the growth surface. Sodium oleate shows better adhesion to the () facet, leading to the growth of nanocubes

Nanorods and wires

Plasma sputtering CH and N plasma sputtering of hematite () wafers; energetic stability of the () faces promotes one-dimensional growth along the [] direction [] Chemical vapour deposition

[Fe(OBut)] precursor deposited layer by layer on alumina and MgO substrates; gold applied as a catalyst at  °C, resulting in a preferential growth along the 〈〉 direction []

Nanorings

Hydrothermal synthesis

Production of hematite nanorings by changing the ratio between phosphate and sulphate stabilising ions; these ions present strong adhesion to the () and () planes and weak affinity to the () plane, causing preferential growth along the 〈〉 direction []

Nanoprisms

Hydrothermal synthesis

Iron(III) chloride applied as a precursor, ethylene glycol as a solvent, ,-propanediamine as a surfactant, and sodium acetate as an additive; –NH group in ,-propanediamine coordinates Fe+ ions, changing the surface conditions of the FeO crystals. Nanoprism morphology is controlled by changing the ratio between ethylene glycol and ,-propanediamine []

synthesis and thermal decomposition, deeply influencing the particle size and morphology [2]. The most important factor for controlling the shape of the prepared nanocrystals is tuning the growth rate of the specific facets. For instance, octahedral nanocrystals are surrounded by eight {111} planes, which are generated when the growth rate along the [100] direction is faster than that along the [111] direction [2, 37]. In contrast, nanocubes are obtained when the growth rate along the [111] direction is faster than that along the

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[100] direction because the particles are bounded by the {100} planes [2, 37]. Zhou et al. [38] succeeded in controlling the shape and surface structure of iron oxide nanocrystals synthesised by the thermal decomposition of iron oleate in the presence of sodium oleate (Figure 3.8). In this case, sodium oleate preferentially binds to the Fe3O4 {111} facets under mild conditions with the application of 1-octadecene as a solvent. This procedure results in the crystallisation of the {111} facet of the exposed iron oxide plates, yielding truncated octahedrons and tetrahedrons (Figure 3.8). In contrast, cubic, concave, multibranch, and assembled Fe3O4 structures with {100} facets are obtained by varying the molar ratio between iron and sodium oleates, applying high boiling temperatures, and using tri-noctylamine as a solvent [38].

Figure 3.8: High-resolution transmission electron microscope (HRTEM) and TEM images of iron oxide (a) and (b) plates, (c)–(f) truncated octahedrons, and (g)–(i) tetrahedrons (reprinted with permission from reference [38], Copyright 2015 American Chemical Society).

3.4 Advanced studies on iron oxide formation As discussed in detail in Part II of this book, studying the formation of the structure of solid materials, such as iron oxide, is critical for understanding and optimising the synthesis strategies, in addition to controlling the structure and structure-related

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properties by influencing the reaction conditions. Examples of these properties relevant to iron oxide particles include crystal size, size distribution, morphology, and magnetism. Interestingly, the crystallisation process of iron oxides involves the formation of pre-nucleation clusters (PNC), deviating from the classical nucleation theories [39]. During the formation of iron(III) oxide, solutions of iron(III) ions are prepared, in which their hydrolysis initially leads to the formation of iron hydroxide species. Low pH values favour the stabilisation of mono- and dinuclear species, which polymerise when the pH value increases above approximately 6, resulting in the formation of oxo-bridged poly-iron complexes [39]. These hydrolysed iron(III) solutions comprise species with low sedimentation coefficients, identified as 2–4 nm spheres by electron microscopy, posteriorly agglomerating to form rods upon ageing, which transform into rafts and flocculate. Hydrated amorphous iron(III) hydroxide species are then formed and consecutively converted into α-FeO(OH) and α-Fe2O3 through dehydration and crystallisation [39]. According to Sun et al. [40], the non-classical aggregation-based pathway for crystallisation also explains the heat-induced crystallisation of magnetite. Therefore, the domains of the pre-aligned clusters consist of a direct precursor for the iron oxide crystals. The application of the amphiphilic ligand, 10,12pentacosadiynoic acid (PCDA), enhances the mobility of the clusters and, therefore, their crystallisation [40].

3.5 Applications

Figure 3.9: Synthesis mechanism and applications of Fe3O4@SiO2@APG-F nanocomposites (republished with permission of Royal Society of Chemistry from reference [41]; permission conveyed through Copyright Clearance Center, Inc.).

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Owing to their low production cost, high physical and chemical stabilities, biocompatibility, environmental safety, and large surface area [13], iron oxide nanoparticles are widely used in diverse biomedical and technological fields, such as MRI, cancer treatment, drug delivery, bioimaging, and magnetic separation, in addition to their application as energy materials and nanocatalysts. MRI is a powerful, non-invasive diagnostic technique with elevated spatial resolution based on the detection of proton relaxation in an external magnetic field. The signal contrast can be enhanced by the application of contrast agents [42]. Using MNPs such as Fe3O4 as contrast agents, a magnetic moment is induced on these materials by applying a magnetic field B0, perturbing the magnetic relaxation processes of the water protons around them. This process causes a reduction in the spin–spin relaxation time and darkening of the corresponding area in the magnetic resonance images [2]. For cancer treatment, MNPs are frequently used in hyperthermia therapy. In this case, the MNPs absorb the alternating current energy and convert it into heat. Consequently, the temperature increases to the range 41–46 °C and causes local overheating of cancer tissue, prompting the cancer cells to undergo apoptosis [42]. Combining magnetic and luminescent properties allows the biomedical applications of these compounds to be extended from MRI and hyperthermia to fluorescence-based diagnostic bioimaging [19]. Drug delivery is a pharmaceutic approach for transporting medicaments to a desired location in the body by means of intravenous delivery, which aims to reduce the toxic dose for non-cancerous cells and increase the therapeutic effect at the targeted locations. MNPs are ideal carriers for this application because they can be coated with active tumour-targeting agents. Generally, drug release is triggered by external stimuli, for example, a local temperature increase (hyperthermia) or low pH value, which is typical for cancerous tissues [42]. The magnetic field is applied to accumulate Fe3O4 particles around the target tissue [2]. Moorthy et al. [41] recently produced Fe3O4@SiO2@APG-F multifunctional nanostructured composites, prepared by coating Fe3O4@SiO2 nanoparticles with amine-polyglycidol (APG) and fluorescein isothiocyanate (FITC), for application as hyperthermia agents, fluorescent contrast agents for bioimaging, and drug delivery carriers (Figure 3.9) [41]. Magnetic separation of ions, pollutants, cells, proteins, antibodies, and other bioactive substances is advantageous, in comparison to conventional separation methods such as chromatography, because it has high versatility, is fast and cost-effective, and does not require pretreatment of the active material [42]. According to Franzreb et al. [43], the requirements of magnetic sorbents are a large number of functional groups per mass of magnetic material and saturation of the magnetisation value [43]. Moreover, the high binding proclivity for the target substances and fast magnetic response of the composite should be considered [42]. MNPs such as Fe3O4 can additionally be used as nanocatalysts [2] for the photocatalytic removal of organic pollutants from aqueous solutions. In photocatalysis, light energy is used to conduct oxidation and reduction reactions. In such photocatalytic oxidation reactions, the photocatalyst absorbs light, exciting electrons from the valence to the conduction band, in which the resulting

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electron–hole pair reacts further with oxygen and water to form peroxide anions and hydroxyl radicals, which can oxidise the organic pollutants. The disadvantage of conventional photocatalysts, such as TiO2 and ZnO, is the large band gap energy, which requires high-energy photons (UV), yielding a rather low efficiency with visible light. Iron oxide nanoparticles can offer a solution for this problem owing to their narrow bandgap energy, which combined with their high surface area and magnetisation, enables easy magnetic separation and recycling of the catalyst [42]. When applying MNPs for the photocatalytic degradation of pollutants, toxicity [44] is particularly important because the particles might have direct contact with humans through skin contact, inhalation, or ingestion through drinking water. As a final example, iron oxide nanocrystals, such as Fe3O4 [2] and α-Fe2O3 [25], are attractive materials for the production of negative electrodes for lithium and sodium batteries, particularly because of their low cost, high stability, non-toxicity, low environmental impact, and high corrosion resistance (α-Fe2O3) [25].

3.6 Surface functionalisation For the application of magnetic nanocrystals in most of the fields mentioned in Section 3.5, additional surface functionalisation is necessary. In particular, this process is relevant for three main reasons: (i) to increase the chemical stability, (ii) to combine multiple properties, and (iii) to improve water dispersibility. Owing to their high surface area, magnetite nanoparticles can easily oxidise to iron(III) oxide in air [2]. Consequently, after a short time, γ-Fe2O3 can become the major phase on small particles (for example, diameter of 5 nm). Therefore, the particles become less sensitive to oxidation with increasing crystal size [2]. In addition to avoiding the loss of magnetisation [13], increasing the chemical stability by surface functionalisation is essential for biomedical applications to prevent the degradation of the nanoparticles in vivo, for example, caused by lysosomal enzymes, and therefore prevents the release of iron cations in the body, which could cause severe oxidative reactions in cells [42]. The functionalisation of iron oxide nanocrystals allows the combination of their intrinsic magnetism with supplementary properties, such as biocompatibility, luminescence, or mesoporosity, by introducing interactive functions on the MNP surface through coating, for example, with polymers, biomolecules, silica, metals, luminescent materials, or even metal–organic frameworks [13, 42]. Improving the water dispersibility is required for biomedical applications [13, 42]. Bare MNPs are typically prone to aggregation, owing to their large surface-to-volume ratio and high magnetisation [45]. Functionalisation influences the hydrodynamic diameter and surface charge, improving the colloidal stability. Coating the surface with silica by the so-called Störber method [47], in which tetraethyl orthosilicate (TEOS) or its derivatives are applied as precursors, is the most common functionalisation procedure for iron oxide nanoparticles. In addition to

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Figure 3.10: TEM images of Fe3O4@SiO2 nanoparticles produced with (a) 75 µL, (b) 150 µL, (c) 300 µL, and (d) 600 µL TEOS, resulting in different SiO2 shell thicknesses (reproduced with permission from reference [47]. Copyright 2012 American Chemical Society).

increasing the biocompatibility, this process enhances the chemical and mechanical stability against variations in pH and temperature, reducing the oxidation of Fe3O4 nanomaterials to Fe2O3 at high temperatures [45]. A silica coating is an excellent surface modifier, owing to its non-toxicity and ability to conjugate additional functional groups for coupling and labelling biotargets. For practical applications, every iron oxide nanocrystal in a batch must be coated with a homogeneous SiO2 shell because an unequal core number, the formation of core-free SiO2 particles, or an inhomogeneous thickness of the SiO2 shell results in an irregular magnetic field and a broad particle size distribution. Ding et al. [46] systematically studied the coating process during the production of Fe3O4 nanoparticles using a reverse microemulsion approach (Section 3.3). For this purpose, the surfactant Igepal® CO-520 was dispersed in cyclohexane by sonication. An Fe3O4 suspension in cyclohexane was added to this solution, followed by an ammonium hydroxide solution and different amounts of TEOS (Figure 3.10). The resulting core–shell nanoparticles were centrifuged, washed, and redispersed in ethanol. The coating procedure occurred via a ligand-exchange process involving a surfactant, an oil phase, and hydrolysed TEOS. The thickness of the SiO2 shell could be adjusted by controlling the quantities of Fe3O4 nanoparticles and TEOS [46]. As discussed by Hola et al. [42], three main coupling strategies are used to enhance the compatibility of MNPs to medical applications. These strategies are based on the chemical modification of amino groups and bioactive substances on the nanoparticle surface and activation of carboxyl groups for enzyme immobilisation, including covalent and non-covalent approaches. Some functionalisation examples include (i) polymer-grafted nanoparticles, (ii) nanoparticles grafted with low-molecularweight compounds, (iii) nanoparticles coated with bifunctional polymers (PEG),

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Figure 3.11: Common strategies for functionalising magnetic nanoparticles (derived from reference [42]).

(iv) stimuli-responsive hydro/nanogels, (v) magnetic clusters, (vi) porous polymer beads, (vii) micelles, (viii) inorganic shells, and inorganic nanohybrids with (ix) magnetic cores or (x) magnetic shells (Figure 3.11). For MRI applications (Section 3.5), low-molecularweight compounds enhance water relaxivity, while functionalisation with hydrophilic polymers, such as PEG and dextran, avoids fast opsonisation of nanoparticles because they do not strongly interact with opsonin proteins, such as albumins or immunoglobulins [42]. In contrast, surface engineering with polymeric materials containing functional groups that are responsive to coupling allows the immobilisation of antibodies and glycoproteins for passive targeting of cancer cells, improving early diagnostics. In addition, to attach drugs to MNPs, drug molecules can be loaded in the polymer interspaces of magnetic clusters or encapsulated in stimuli-responsive hydrogel/polymer frameworks [42]. MNPs, encapsulated by a porous polymer-bead matrix, are used in chemical separation-related applications and core–shell approaches with polymeric or inorganic shells, resulting in an increased surface area, enabling the incorporation of a large number of functional groups for chemical separation. Inorganic shells, such as Ag-coated MNPs, provide antibacterial, antifungal, and antiviral properties, which can be used in antimicrobial targeting [42]. Finally, the functionalisation of magnetic cores with inorganic and hybrid shells is explained in more detail in the following sections. The disadvantages of functionalisation are typically related to the increase in size due to the addition of supplementary coating layers, resulting in composites above the nanometre range and, often, also reducing the magnetisation.

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3.6.1 Coating with luminescent materials Bifunctional materials, combining optical and magnetic properties into a single nanostructured entity, are advantageous because of their easy manipulation via a magnetic field and diverse chemical functionalities. Bifunctional magnetic–luminescent nanomaterials can potentially be applied in in vitro and in vivo labelling, multifunctional biomarkers, photothermal destruction of tumour cells, and as reusable catalysts that can be separated by a magnetic field, or in pre-concentration procedures for quantitative DNA analysis [18, 45]. The general preparation strategies for these bifunctional materials include, for example, (i) coating or layer-by-layer deposition of lanthanide ions or quantum dots on magnetic-core nanoparticles, (ii) encapsulation of magnetic particles within luminophore-containing polymers, and (iii) functionalisation of MNPs with fluorescent dyes or d-transition metal complexes [18]. Lanthanide (Ln) ions [48–51] are remarkable candidates for supplying magnetic particles with additional luminescent properties. For example, trivalent lanthanide ions are characterised by their well-defined narrow emission peaks in spectral ranges varying from visible to near infrared, with long lifetimes and high quantum yields. Most trivalent lanthanide ions are paramagnetic because of the unpaired electrons. Their magnetic properties are usually determined by the ground state configuration because the excited and the ground states are well separated, owing to the spin–orbit coupling, and are thermally inaccessible. The magnetic moments of trivalent lanthanide ions can contribute to the total magnetisation of these magnetic–luminescent nanomaterials [44]. However, iron oxide is a strong luminescence quencher, owing to the killer effect [52] caused by the energy transfer process that occurs when luminescent materials such as Ln ions are in direct contact or close to the metal oxide surface. To overcome this problem, an intermediate layer or energy barrier must be introduced between the iron oxide surface and the luminescent layer. Typically, applied protective materials include polymers, silica, or semiconductors [18]. For the selective adsorption of transition metals and trivalent lanthanide ions, for example, from aqueous solutions, the surface of the magnetic particles can be functionalised with ethylenediaminetetraacetic acid (EDTA)-based groups. Additionally, attaching lanthanide-based complexes to these composite EDTA-covered magnetic particles qualify them for magnetic–luminescent applications, in which lanthanide complexes with carboxylates, heteroaromatics, and β-diketonate ligands are particularly advantageous, owing to the so-called antenna effect [53]. Owing to this effect, these lanthanidebased complexes present strong energy absorption by the ligands and efficient energy transfer to the light-emitting lanthanide ions. The ligands can also protect the lanthanide ions against high-energy oscillator groups such as OH and NH, decreasing their quenching effects. The chelating effect of the hexadentate EDTA is advantageous for the anchoring of metal complexes on the particle surface to avoid leaching of the metal ions [45, 48]. For example, Pires et al. [45] functionalised magnetite particles with

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organosilanes containing EDTA derivatives attached to Eu3+ β-diketonate complexes to decrease the quenching effect of Fe3O4 on the lanthanide luminescence. Khan et al. [48] applied a similar core–shell approach using a Fe3O4 core coated with a SiO2 protective shell, and further grafted with Eu3+ and Tb3+ complexes, resulting in Fe3O4@SiO2-(TTALn-L) materials, in which Ln is Eu3+ and Tb3+ and L represent different ligands, such as thenoyltrifluoroacetonate (TTA), thioglycolate (TC), 4-aminobenzoate (AB), and 4-(aminomethyl)benzoate (AMB). In addition to increasing the stability (as described above), the SiO2 layer can be used to embed Ln3+ complexes [45]. Another example of magnetic–luminescent composites is triple-doped Fe3O4/ ZnS@LaF3:xCe3+,xGd3+,yTb3+ (x = 5; y = 5, 10, and 15 mol%) nanomaterials [18]. To prepare these particles, Fe3O4 nanoparticles were functionalised with ZnS as an energy barrier between the magnetic core and the luminescent shell, and coated with LaF3: Ln3+ via a chitosan-assisted co-precipitation method. In this case, coating with ZnS presents two main advantages. In addition to being luminescent and a classical scintillating material, ZnS is a semiconductor and, in contrast to the isolating SiO2 protective layers, might not severely interfere with the magnetic properties of the iron oxide core. Regarding the luminescent properties, a broad blue emission (400–550 nm) was observed under UV irradiation, which was assigned to the vacancies on the Fe3O4/ZnS surface. Moreover, the respective emission spectra also presented narrow emission lines due to the 5D4 → 7F6-0 Tb3+ transitions, resulting from the Ce3+ → Gd3+ → Tb3+ energy transfer. These composites have potential applications in the production of magnetic light-converting molecule devices (MLMCDs) and high-energy detection [18].

Figure 3.12: Example of an inorganic functionalisation shell combining the properties of a magnetic Fe3O4 core (in black) with an afterglow SrAl2O4:Eu2+,Dy3+ shell (in green), separated by a SiO2 protective layer (in grey) (derived from reference [54] with permission from John Wiley and Sons).

As exemplified above, most combinations of magnetic and optical properties in single assemblies consist essentially of a magnetic core, coated with a protective middle layer and a fluorescent shell, generally comprising small luminescent molecules, complexes, or quantum dots [54]. The disadvantage of these materials for application as

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biomarkers is their ability to emit light only simultaneously with the excitation process. Therefore, they are affected by the auto-fluorescence problem, in which the fluorescent materials within living cells are excited simultaneously with the marker during bioimaging [54]. A solution to this problem is the production of multifunctional assays, which combine magnetic and afterglow properties. Hence, the marker can emit light after the excitation source is blocked, achieving a higher contrast with the target cells. Nevertheless, the afterglow process is achieved by more challenging mechanisms than fluorescence and requires, for instance, the construction of an additional crystalline layer on the magnetic particle surface, profiting from the combination of lattice defects and energy transfer between mixed-valence lanthanide ions such as SrAl2O4:Eu2+,Dy3+. Recently, the production of the first magnetic afterglow particles [54] was achieved by incorporating a combustion synthesis approach [50] into a double core–shell strategy for the production of Fe3O4@SiO2@SrAl2O4:Eu2+,Dy3+ nanoparticles (Figure 3.12). Example applications for these nanoparticles include the identification of cancer cells, detection of relevant biomolecules, and separation of cells in vitro [54]. Figure 3.13 shows the high magnetisability of the magnetite core particles, which were easily separated from the liquid media. The resulting Fe3O4@SiO2@SrAl2O4:Eu2+,Dy3+ nanoparticles presented a broad emission band in the green spectral range with a maximum at 520 nm, assigned to the 4f65d1 → 4f7 (8S7/2) transition of Eu2+ in the SrAl2O4 host lattice [50], were able to emit light for a few seconds after the excitation source was turned off, and could be attracted by a magnetic field (Figure 3.13) [54].

Figure 3.13: Magnetic and afterglow properties of Fe3O4@SiO2@SrAl2O4:Eu2+,Dy3+ nanoparticles (derived from reference [54] with permission from John Wiley and Sons).

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3.6.2 Coating with metal–organic frameworks (MOFs) As explained in detail in Chapters 2 and 6, MOFs are a new class of highly ordered porous materials that have emerged in the past two decades, and are composed of inorganic units, for example, metal ions and clusters that are interconnected by functional organic linkers, resulting in a porous framework [55]. These materials are advantageous because of their elevated surface area, tailorable pores, high porosity, distinct host–guest interactions, typical intrinsic light-emitting properties, and good thermal stability. Hence, these properties justify the wide application of MOFs in gas adsorption and storage, drug delivery systems, luminescence, chromatography, and sensing [53, 55]. Interestingly, MOFs can also be combined with MNPs to produce multifunctional core–shell composites, maintaining the functions of both constituents in a single material. Wehner et al. [55] joined several functionalisation approaches by combining iron oxide particles with luminescent MOFs. In this work [55], Fe3O4/SiO2 composites
 were coated with the 2∞ Ln2 Cl6 ðbipyÞ3 · 2bipy (bipy, 4,4′-bipyridine) MOF, allowing the magnetic collection of these bifunctional particles from fluids and their concentration for better luminescence detection. Owing to their defined porosity, luminescent MOFs can be applied as sensors for small molecules, as demonstrated by the
 Fe3 O4 =SiO2 = 2∞ ln2 Cl6 ðbipyÞ3 · 2bipy composite, which influences the luminescence properties of chemical compounds through interactions such as quenching. This enables the use of these materials in sensing applications, for example, detecting water traces in organic solvents [55]. Another example of combining MOFs with magnetic particles was demonstrated by Quin et al. [56], who prepared carboxyl-functionalised magnetic cores, homogeneously coated with the zirconium-based UiO-67 (UiO = University of Oslo) MOF through liquid-phase epitaxy (LPE). These core–shell particles were applied as the stationary phase in high-performance liquid chromatography (HPLC) columns with different phenol derivatives, with acetonitrile/water as the mobile phase. MOFs are attractive materials for preparative chromatographic separation because of the absence of dead volume, good solvent stability, and high loading capacity [56]. However, the heterogeneous particle size and morphology of MOFs are limiting factors for obtaining quantitative kinetic chromatography data. This problem is solved by coating surface-attached MOFs (SURMOFs) on well-defined magnetic particles, combining LPE with a core–shell approach. SURMOFs are also advantageous because of their controllable thickness. In this case, the magnetic properties are important for the frequent solid–liquid separation steps performed during the synthesis of the MOF shell by LPE. The uptake kinetics are dominated by intraparticle diffusion, and the pore diffusivities of the phenol derivatives are approximately 1.3 × 10–13 m2/s [56]. Another MOF, MIL-100(Fe) (MIL = Materials of Institute Lavoisier), has been applied to coat the surface of magnetic particles via LPE and tested for chromatographic applications (Figure 3.14) [57]. The rigid 3D structure of MIL-100(Fe) is advantageous because of its

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high stability in aqueous media, large pores (2.5–2.9 nm), low toxicity, and easy accessibility of the iron metal ions [57].

rp

rc

Structure of MIL-100(Fe)

Figure 3.14: Schematic of magnetic cores coated with a MIL-100(Fe) shell (reprinted from reference [57], with permission from Elsevier).

3.7 Conclusions and outlook This chapter offers a quick overview on the vast literature on the synthesis, functionalisation, and applications of magnetic nanoparticles, with a focus on iron oxidebased materials. In addition to their high magnetisation, iron oxide-based nanoparticles are advantageous, owing to their low cost, and can be produced by a variety of methods, such as co-precipitation, thermal decomposition, hydrothermal and solvothermal synthesis, sol–gel synthesis, the polyol method, microemulsion, sonochemical synthesis, microwave-assisted synthesis, and biosynthesis. These different synthesis approaches enable the flexible production of nanocrystals with several different morphologies, such as cubic particles, octahedral particles, nanowires, and nanorods. The production of these magnetic nanomaterials is particularly important for their applications in the biomedical and technological fields, such as MRI, cancer treatment, drug delivery, bioimaging, and magnetic separation, in addition to their applications as energy materials and nanocatalysts. To achieve this objective, iron oxide nanomaterials must be additionally functionalised by coating with polymers, biomolecules, silica, metals, luminescent materials, or metal–organic frameworks to increase their chemical stability, combine multiple properties with intrinsic magnetism, and improve their water dispersibility. Despite the immense advances in the development of magnetic nanomaterials in the past few years, several challenges remain to be overcome. It is difficult to achieve absolute control over the shape and size distribution [13] using many of the existing synthetic and functionalisation approaches, and additional toxicological research is required [13]. Concerning toxicological issues, supplementary effort is required to define standard criteria to evaluate toxicity to improve the comparability between different

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studies [13]. The same problem of developing standard criteria for improving the comparability between experiments is faced for evaluating the photocatalytic properties [44] of iron oxide nanoparticles, which must be addressed. Additional current challenges lie in the optimisation of the targeting efficiency and long-term stability in in vitro and in vivo studies [13], as well as the preservation of the high magnetisation and nanoscale dimensions, with the application of thick isolating functionalising coating layers. Among the future trends in the functionalisation of magnetic nanomaterials is the further development of multifunctional layers, such as applying fluorophores and radiotracers [13], or biomolecules for the application of magnetic nanoparticles in theranostic nanomedicine. To tailor structure-related properties such as magnetism, particle size, and morphology, additional information about the formation mechanism of these materials is required, which therefore requires further application of in situ characterisation techniques [39, 40] during their synthesis. Finally, the development of methods for the large-scale production of (functionalised) iron oxide nanoparticles, in addition to the further development of MOF coatings for magnetic materials, should also be the focus of future investigations within this exciting research field.

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Niklas Ruser✶, Huayna Terraschke

Chapter 4 Nanostructured bioceramics Bioceramics are a very promising research field, especially for tissue engineering. Their biocompatibility with the human body enables several applications, and their beneficial properties can be additionally improved by scaling the micro- to the nanostructure. This chapter presents an overview of different nanostructured bioceramics, how they can be synthesised and functionalised, typical applications, and a brief outlook for this research topic.

4.1 Introduction Ceramics are crystalline inorganic solids generally manufactured at high temperature and distinguished by their thermal and chemical stability. The exact properties of any ceramic depend, besides on their composition, on its crystalline structure, which is determined for instance by the synthesis method [1]. Historically, ceramics have been used for medical applications owing to their chemical stability. Today, a more sophisticated understanding of ceramics has enabled bioceramics to be produced. Rather than being fully inert in the human body, bioceramics undergo targeted, beneficial reactions with specific tissues to facilitate medical procedures and/or healing [2]. Bioceramics can be sub-categorised as bioinert or bioactive, depending on their biochemical and physical properties and whether they interact with the surrounding tissue. In addition, each bioceramic can be classified as either “non-resorbable” or “resorbable” according to its macrostructure (e.g. porous, dense, powdered, granulated, or coated films) [2]. These materials are mostly prepared from, calcium phosphates, alumina, zirconia, silica-containing substances, and other compounds. Calcium-phosphatecontaining bioceramics perform well in medical applications because of their similarity to inorganic bone scaffolds [2]. This chapter focuses on calcium-phosphate-, calciumcarbonate-, and glass-based materials. In the 1950s, the condition for tissue repair using inert ceramic compounds was that ceramics should be biocompatible; they should not interact with living tissue but should replace damaged ones. The intention was to use non-toxic and non-biodegradable materials so that the organism was not harmed. Such materials are identified as foreign substances by the immune system and elicit a simple encapsulation response. The ✶

Corresponding author: Niklas Ruser, E-Mail: [email protected]

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discovery of bioceramics in the 1980s, however, triggered a change in thinking about the purpose and intention of ceramic bone treatments. The main reason for the application of calcium-phosphate-based ceramics is that they exhibit mineral phases very similar to those of bones and teeth. Therefore, these second-generation bioceramics do not replace damaged tissue but repair it owing to their bioactivity and biodegradability. Recently, research in this field has progressed towards third-generation bioceramics by adding substances that can be continuously released from the applied materials into the host tissue, enabling biologically active compounds such as hormones, growth factors, and tailored drugs to react at a specific target zone to regenerate damaged tissue. Additionally, hierarchical bone porosity has been mimicked to act as a cellular scaffold (Figure 4.1) [2].

Figure 4.1: Scheme showing evolution of bioceramics [3].

In general, it is important to differentiate between nanobiomaterials and nanostructured biomaterials. Nanobiomaterials include single molecular compounds like proteins. Other biomaterials consisting of multiple molecules forming clusters or crystallites within the range 1–100 nm are the nanostructured ones. In particular, nanostructured bioceramics exhibit many unique properties that make them better candidates than microstructured bioceramics for e.g. medical or dental applications [4]. For example, the surface area-to-volume ratio of nanoparticles results in new functionalities, explained in detail in the following chapter [5].

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4.2 Types of nanostructured bioceramics 4.2.1 Unique properties of medically relevant nanobioceramics So far, various nanostructured bioceramics have been synthesised ranging from simple ceramics to porous materials [2], both of which exhibit several unique properties that make them exceptionally well-suited for medical application in bone and tooth repair. As defined by the European Commission, materials are “nanomaterials” when at least one of their dimensions is smaller than 100 nm. Nanomaterials offer great advantages compared to their micromaterial counterparts, especially because nanomaterials exhibit a high surface area-to-volume ratio (Figure 4.2) and quantum effects. Although quantum mechanical effects are not applicable to biomaterials, the surface area-to-volume ratio strongly influences bioceramic properties [4]. For comparison, the percentage of atoms on a microparticle surface accounts for approximately 1 % of all the particle atoms. If the particle diameter is decreased to the nanometre range (10 nm for example) the proportion of atoms on a microparticle surface increases to approximately 10 %. For instance, a ⌀2 nm nanoparticle exhibits over half of its atoms (~60 %) on the surface [4].

Figure 4.2: Cutting particle without losing any mass: volume remains the same while surface area increases. Cut surfaces (blue) are additionally created surface areas.

In summary, there are five main reasons for the improved applicability of nanobioceramics compared to their microbioceramic counterparts. First, for the same particle volume, nanobioceramics exhibit more surface atoms than microbioceramics, which means that far more atoms are available for bonding to tissue surfaces such as bone [5]. Second, because bioceramic nanoparticles are smaller than their bioceramic microparticle counterparts, the former can easier diffuse into organic tissue. Third, nanobioceramics are highly sinterable, making them easier to process and handle even though they self-aggregate. Fourth, the development of nanomaterials enables chemical homogeneity and structural consistency, which is a great achievement for producing biomaterials that imitate hard tissues [4]. Fifth, healthy bone matrix

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consists of crystalline nanostructures. Therefore, using comparable structures in hardtissue repair eases the integration of bioceramics into healthy bone structures [5]. In addition, medically used nanobioceramics can exhibit high porosity [4] and bone exchangeability [6]. Hard-tissue repair is then more effective when highly porous bioceramics are used because porous surfaces exhibit exceptional adsorption properties. This enables efficient protein adsorption, which is a key step in bone repair [4]. Hence, the bone exchangeability of nanobioceramics enables implants to be fully replaced by healthy, natural bone tissue. To accomplish this, the bioceramic must be osteoconductive (i.e. support bone tissue growth in and around implant structures), bioresorbable (i.e. enable the host organism to break down implant structures) [6], and biodegradable. The following sections outline the most important nanobioceramics that best meet these extensive criteria and are either already regularly applied or are currently being developed for application to medical treatments.

4.2.2 Calcium-phosphate-based bioceramics Calcium phosphates consist of a large family of different compounds varying by the calcium-to-phosphorus ratio, additional anions (e.g. OH− and F−), or hydration level (Table 4.1). By weight, natural bone consists of 69 % calcium phosphate crystals [8], providing the tissue with the necessary stiffness and compression stability [2]. The other components are collagen (20 wt.%), water (9 wt.%), and traces of various other organic compounds, making calcium phosphate ideal for bone repair [8]. Calcium phosphate bioceramics are also bioactive, non-toxic [2], lightweight, chemically stable, and resistant to changes in the surrounding physiological fluid and microbes [8]. Calcium phosphate nanoparticles can exhibit various morphologies, as shown in Figure 4.3 [5]. Despite their numerous advantageous properties, conventional calcium-phosphatebased microbioceramics still present some technical challenges. First, their relatively small surface areas (2–5 m2/g) result in poor sinterability and low mechanical strength [8]. Additionally, clinical applications are limited because calcium-phosphate-based microbioceramics do not completely bond to intact hard tissue, resulting in poor durability for long-term bone implants [10]. However, many of these issues can be overcome by reducing bioceramic microparticles to nanoparticles [5] and increasing the specific surface area to more than 100 m2/g [4]. Commercially available calcium-phosphatebased bioceramics for medical applications include hydroxyapatite (HAp), α/β-tricalcium phosphate (α/β-TCP), bi-phasic calcium phosphate (BCP), monocalcium phosphate monohydrate (MCPM), and non-sintered apatite (AP) [8].

Table 4.1: Major calcium phosphate species exhibiting Ca/P ratios between 0.5 and 2.0 [4]. Chemical formula

Ca/P ionic pH stability range in aqueous ratio solutions ( °C)

Representative size of synthesised nanoparticles (nm)

Monocalcium phosphate monohydrate (MCPM)

CaðH2 PO4 Þ2 · H2 O

.

.–.

d

Anhydrous monocalcium phosphate (AMCP)

CaðH2 PO4 Þ2

.

b

d

Dicalcium phosphate dihydrate CaðHPO4 Þ · 2H2 O (DCPD)

.

.–.

d

Anhydrous dicalcium phosphate (ADCP)

CaHPO4

.

b

–

Octacalcium phosphate (OCP)

Ca8 ðHPO4 Þ2 ðPO4 Þ4 · 5H2 O

.

.–.

d

.

a

~– [] / ~

.–.

~– c



α/β-Tricalcium phosphate (α/β- α =β−Ca3 ðPO4 Þ2 TCP) Amorphous calcium phosphates (ACPs)

Cax Hy ðPO4 Þz · nH2 O ðn = 3 − 4.5, 15 − 20 wt% H2 OÞ

Calcium-deficient hydroxyapatite (CDHA)

Ca10 − x ðHPO4 Þx ðPO4 Þ6 − x ðOHÞ2 − x ð0 < x < 1Þ .–.

.–.

~

Hydroxyapatite (HAp or HA)

Ca10 ðPO4 Þ6 ðOHÞ2

.–

– []

.

(continued)

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Table 4.1 (continued) Chemical formula

Ca/P ionic pH stability range in aqueous ratio solutions ( °C)

Representative size of synthesised nanoparticles (nm)

Fluorapatite (FA)

Ca10 ðPO4 Þ6 F2

.

–

– []

Oxyapatite (OA, OAp, or OXA)

Ca10 ðPO4 Þ6 O

.

a

d

.

a

d

Tetracalcium phosphate (TTCP) Ca4 ðPO4 Þ2 O a) Compounds that cannot be precipitated from aqueous solutions. b) Only stable above 100 °C. c) Always metastable. d) No example was found in the available literature.

Niklas Ruser, Huayna Terraschke

Compound

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Hydroxyapatite (HAp) Hydroxyapatite (HAp; Ca10(PO4)6(OH)2) [11] is a calcium-phosphate-based bioceramic that forms plate-like or needle-shaped nanocrystals [8] and is the foundational mineral component of bone tissue, wherein HAp crystals can, for instance, range from 30–50 nm long × 15–30 nm wide × 2–10 nm thick [5] to 40–60 nm long × 20 nm wide × 1.5–5 nm thick [8]. Tooth enamel is also primarily composed of HAp crystals, typically 100 µm long × 25 nm wide × 70 nm thick [11]. Despite their larger size, sintered HAp microparticles are decidedly well-suited for application to both hard and soft tissues (e.g. bone, skin, and muscle) because HAp microparticles exhibit slow biodegradability, osteoconductivity, osteoinductivity [12], and moderate surface roughness (e.g. approximately 10 nm for crystals measuring ~ 180 nm across). Nanocrystalline HAp further improves many of these properties. The smaller crystals exhibit higher surface roughness (approximately 17 nm for crystals ~ 67 nm across) making them more adept at molecular adsorption or surface adhesion (Figure 4.4) [4]. Additionally, nanocrystal bioactivity is higher than that of conventional crystals [12]. Although neither micro- nor nanocrystalline HAp exhibits particularly high mechanical strength (which limits their overall application in load-bearing hard-tissue repairs) [12], nanocrystalline HAp exhibits greater hardness and toughness [4]. The improved adhesion provided by larger contact surfaces makes nanocrystalline HAp more sinterable, thereby improving the ceramic density and (theoretically) fracture resistance [8, 12]. Vertebrate bones are based on the biological apatite mineral phase. The term “apatite” describes crystalline solids exhibiting the general formula M10(ZO4)6X2, wherein ions in the crystal structure can be replaced by other ions (Table 4.2). Such hybrid crystals more accurately reflect the composition of healthy bone tissue because its hydroxyapatite (Ca10(PO4)6(OH)2)-based structure exhibits many exchanged ions. In bone apatite, the main ions replacing Ca2+ are Na+, Mg2+, and K+ and replacing PO43− and OH− are Cl−, F−, and mostly CO32− [11]. As the HAp-based bone crystal structure contains a mixture of different ions, a more precise chemical formula for bone apatite would be (Ca,Mg,Na)10(PO4,CO3)6(OH)2 [13]. Table 4.2: For apatite chemical formula M10(ZO4)6X2, components can be replaced by various ions [11]. Components

Ions

M ZO X

Ca2+ , Na+ , Mg2+ , K+ , Mn2+ , Zn2+ , Fe2+ , Sr2+ , Ba2+ , Cd2+ , Pb2+ , H+ , Al3+ , etc. PO4 3 − , CO3 2 − , SiO4 4 − , SO4 2 − , VO4 3 − , AsO4 3 − , etc. OH − , CO3 2 − , F − , Cl − , Br − , O2 − , etc.

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Figure 4.3: Morphologies of calcium phosphate nanoparticles: (A) spherical; (B and C) rod-shaped; (D–F) rod shaped/nanowires/tubes; (G–I) nanocrystalline; and (J–L) fibrous-, needle-, and plate-like structures (republished from reference [5] with permission from Royal Society of Chemistry and conveyed through Copyright Clearance Center, Inc.).

Tricalcium phosphate (TCP) Tricalcium phosphate (TCP) exists as α-, αʹ-, and β-TCP polymorphs. For medical applications, however, only the α- and β-isoforms are of interest because αʹ-TCP only exists at excessively high temperatures (>~1,430 °C) below which the αʹ polymorph instantly converts into the α one, which is stable at room temperature. The β polymorph is also stable at room temperature. Although the β structural configuration is formed by heating the α isoform to ~1,125 °C, such extreme temperatures are not required to retain the β configuration. Due to their different structures, α- and β-TCP exhibit

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Figure 4.4: Higher surface roughness increases generated surface area. Thus, rough surfaces provide more space for molecules, ions, or atoms to bond than equivalent area of perfectly flat 2D surface. The top panel depicts a rough surface providing space to hold, as a very simplified example, 12 blue dots, while the bottom panel shows that the equivalent distance across a flat surface can only hold 9. Table 4.3: Densities and solubilities of α- and β-TCP [14]. Compound

α-TCP β-TCP

Density (mg/cm)

. .

Solubility (mg/L)  °C

 °C

. .

. .

different physicochemical properties including density, solubility, and biodegradability, making them well-suited for different medical applications (Table 4.3) [14]. The higher solubility and reactivity of α-TCP make it useful for producing bioceramic calcium-phosphate-based cements. However, some α-TCP granules or blocks are commercially available. Such α-TCP-based cements form solid HAp when exposed to a slightly basic physiological pH (7.2–7.4) because HAp is the more chemically stable form of calcium phosphate under these conditions. β-TCP, on the other hand, is well-suited for producing mono- or biphasic bioactive dense and macroporous granules and blocks [14].

4.2.3 Calcium-carbonate-based bioceramics Calcium carbonate can also be applied to bioceramics and is widely used in industry. Recently, calcium carbonate has also been used in medical applications owing to its biocompatibility, bioresorption, osteoconduction, low cost, safety, and natural abundance. Additionally, calcium carbonate is a candidate for producing controlled-drugrelease systems because of its pH sensitivity, protracted biodegradation, and drug

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loading capability, which enables drugs to be loaded into calcium carbonate particles (either during synthesis or through surface adsorption on fully formed crystals) and then released slowly over an extended period. Alternatively, calcium carbonate may be added to biomedical cements to improve their resorption rates and facilitate cement replacement by bone tissue [6]. Calcium carbonate exists as calcite, aragonite, and vaterite anhydrous crystalline polymorphs, which can be directly synthesised under appropriate reaction conditions, and polymorph stability descends in the sequence calcite > aragonite > vaterite [6]. Trigonal crystalline calcite is the most stable form of calcium carbonate and is one of the most common naturally occurring biomineral phases. Moreover, because of its biocompatibility, biodegradability, and encapsulation and controlled-drug-release capabilities, trigonal crystalline calcite is, among the calcium carbonate polymorphs, widely used in pharmaceutical and biomedical applications [15]. Aragonite exhibits an orthorhombic crystal structure and biocompatibility like that of calcite and is used due to its dissolution, integration, and replacement by bone tissue. Previously, aragonite has been applied to anti-cancer drug carriers and bone repair implants. For example, an aragonite-based macroporous drug carrier developed for application to osteomyelitis treatment exhibited faster resorption than a common HAp/TCP bioceramic [6]. Hexagonal crystalline vaterite is an excellent candidate for drug-release applications. It exhibits favourably high porosity, a large surface area, and dissolution under relatively mild conditions [6].

4.2.4 Glass-based bioceramics Glass-based bioceramics or bioactive glasses (BGs) comprise a unique category of predominantly silica-based biomaterials [16]. Silica-based BG networks are frequently modified by adding Na, Ca, and/or P bonded through non-bridging oxygen atoms to the network [2]. The first silicate-based BG was designed in 1969 by Hench et al. [17] This glass, the renowned 45S5 Bioglass® (SiO2, Na2O, CaO, and P2O5 = 45, 24.5, 24.5, and 6 wt.%, respectively) provides a foundational formula for preparing many currently used BGs. Many modern bioglasses have been synthesised using a modified 45S5 Bioglass® formula to improve the ability of the bioglass to bond to bone tissue. Studies have identified that substituting 5–15 wt.% B2O3 or 12.5 wt.% CaF2 for SiO2 and CaO, respectively, markedly improves tissue bonding. However, BGs prepared using P2O5 contents considerably above 6 wt.% do not bond to bone as well as BGs prepared using 3 wt.% Al2O3 [18]. Although these formulations are the most well-established and widely used, many other formulations incorporate several elements such as fluorine, magnesium, strontium, iron, silver, boron, potassium, and zinc [17]. Even though bioglasses are not like

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bone tissue, they perform well in biomedical applications because of their unique indirect interactions with bone [2]. In detail, upon contact with bodily fluids, bioglasses initiate the formation of a hydroxycarbonate apatite (or carbonated hydroxyapatite) (HCAp) layer, which exhibits exceptionally high surface affinity for living bone tissue. Therefore, BGs that do not exhibit a pronounced capability to bond to living bone tissue are the starting materials for forming the HCAp layer, and network modifiers in the amorphous silica scaffold help form HCAp in physiological fluids [19]. The formation of biocompatible HCAp layers requires several steps (Figure 4.5) as follows. First, sodium (calcium [18] or potassium) cations docked within the silica scaffold are rapidly exchanged for protons in bodily fluids. These protons covalently bond with silica groups on the glass surface to form SiOH moieties. The proton pool depletion in the surrounding fluid raises the pH. Second, the elevated OH− ion concentration causes the scaffold to be nucleophilically attacked and silica to dissolve as silicic acid (Si(OH)4) and additional SiOH groups. Third, depleted in alkaline and alkalineearth ions, the SiOH groups condense to form Si–O–Si bonds, resulting in a 1–2 μmthick amorphous SiO2-rich layer [19]. Fourth, calcium and phosphate ions diffuse from inside the glass scaffold and reach the interface between the silica-rich layer and the calcium- and phosphate-rich bodily fluids to form an amorphous calcium phosphate (ACP) layer. Fifth, the ACP layer crystallises, supported by the migration of dissolved hydroxyl and carbonate (and to a lesser extent, fluoride [18] ions) in the surrounding fluid to form a carbonated hydroxyapatite (HCAp) layer [18, 19]. Once formed, the HCAp layer strongly bonds with the bone mineral phase, securing the implant to the bone tissue [19]. The HCAp layer is then penetrated by newly formed bone tissue while the implant is slowly resorbed [2]. Accepted by the surrounding tissue, the developed HCAp crystals are incorporated into and strengthened by collagen fibres [20]. This gradual replacement of implant structures with healthy bone tissue enables tissue healing, without provoking an inflammatory response in the surrounding tissues or initiating the formation of a fibrous capsule around the implant and damaged tissue [2]. Figure 4.6 shows scanning electron microscopy (SEM) images of HCAp formed on the surface of a glass ceramic submerged in simulated body fluids (SBFs) for 24 h [21]. As their name suggests, SBFs enable the simulation of physiological fluids outside the body (in vitro) so that BG reactivity to actual bodily fluids can be tested prior to introducing BGs into a living organism (in vivo) [20]. Typically, SBFs mimic both the composition and concentration of the inorganic portion of human plasma [20, 22] and may be by prepared by dissolving appropriate amounts of NaCl, KCl, K2HPO4∙3 H2O, MgCl2∙6H2O, CaCl2, and Na2SO4 in distilled water and buffering them to pH 7.4 using tris(hydroxymethyl)aminomethane and HCl [20].

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Figure 4.5: Step 1: Formation of SiOH groups; step 2: dissolution of Si(OH)4; step 3: condensation of SiOH + SiOH; step 4: formation of ACP layer; step 5: crystallisation of HCAp layer.

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Figure 4.6: Formation of HCAp (HCA) on surface of bioactive glass incubated for 24 h in simulated body fluids (SBFs) (republished from references [21, 23] with permission from Wiley, Inc. and conveyed through Copyright Clearance Center, Inc.).

Previous studies have shown that for other nanostructured particles such as titania, alumina, and HAp, the high bioactivity, osteoconductivity, and bacterial toxicity of BGs are likely to be enhanced by scaling micro- to nanoparticles [17], because the building blocks of human tissue are nanostructures (e.g. in bones) and because cells and surfaces interact in nanodimensions [17]. Changes into the nanodimension, the biomaterial surface-to-volume ratio (i.e. specific surface area), defect concentration, surface roughness, hydrophilicity, and wettability all influence cellular interactions. Wettability influences the ability of cells to adhere to the biomaterial, which affects long-term functionality. A low defect concentration and a high surface-to-volume ratio both increase nanoparticle surface energy and bioreactivity. The higher specific surface area compared to microparticles accelerates ion release into the surrounding physiological fluids (solubility) and increases protein adsorption, leading to higher bioactivity. Additionally, nanocrystal surface topographical attributes such as roughness, shape, and size and the ordered or random distribution of the attributes influence cellular behaviour more than the microcrystal counterparts. Evidences of cellular reactions altered by different nanodimensional topographies have been preciously experimentally shown [17]. Another major influencing factor is surface energy. This way, nanoparticles surfaces can be additionally used to manipulate cellular responses to biomaterials. Moreover, the higher surface energy of nanocomposites stimulates protein adsorption, which influences cellular adhesion, spreading, and proliferation, resulting in improved interaction between the implant and the surrounding tissue. Material roughness and surface energy are connected; in fact, the surface energy polar component increases remarkably with surface roughness. Therefore, surface roughness plays also a role in cellular adhesion and proliferation, and the degree of hydrophilicity is also an important factor because it influences biomaterial wettability [17].

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4.2.5 Porous materials Porous materials, which are typically produced from glasses or phosphates [24], can be broadly categorised depending on the average pore diameter. Micro-, meso-, and macroporous materials exhibit pore diameters below 2 nm, ranging from 2 to 50 nm, and above 50 nm, respectively [25]. Furthermore, the pores themselves can be sub-categorised as either ordered or unordered depending on whether they form an organised interconnected pattern advantageous to the adsorption and desorption of specific molecules [24]. Ordered mesoporous materials require the use of templates and surfactants to build a scaffold in which inorganic precursors gather and condense to form a mesoporous structure. Subsequently removing the template generates highly complex and ordered two- or three-dimensional networks of interconnected pores, which can generate a surface area of approximately 1,000 m2/g and pore volumes of up to 1 cm3/g [2]. Mesoporous materials are of special interest because drugs used in drug-delivery applications exhibit an average diameter of one nanometre (Figure 4.7); therefore, they fit well inside pores which exhibit also a large surface area and a narrow diameter distribution [24]. Although mesoporous structures have been shown to improve both in vitro bioactivity and drug delivery capability, bone regeneration is mainly driven by cells. Therefore, mesoporous materials exhibiting pore diameters in the range 2–50 nm are unsuitable for bone cells, which require cavities exhibiting diameters in the micrometre range. For instance, because bone pores exhibit pore diameters in the range 1–3,500 µm, an implant is more appropriate when it exhibits similar macroporosity because bone pores are necessary for physiological functions such as cellular adhesion, penetration, growth, and proliferation leading to bone tissue growth in the implant and subsequent vascularisation (or vasculo-genesis) [26], which is the formation of new blood vessels [27]. Therefore, it is necessary to produce materials exhibiting both mesoporosity to enable drug adsorption and desorption for treating bone pathologies and macroporous structures to sustain the oxygenation, nutrient supply, and vascularisation of bone tissue. This hierarchical porosity has already been manufactured in bioactive glasses. The obtained materials exhibited 3.7 nm-diameter mesostructures, either 200–400- or 500–700 µm-diameter macrostructures, and a bioactive response as the formation of an apatite layer when the material was immersed in SBFs for 4 h [26]. In general, silica-based ordered mesoporous materials have attracted considerable interest for biomedical applications because of their potential in tissue engineering and drug-delivery systems. Such materials exhibit a silica network and a surface composed of silanol groups. As the composition of such materials, especially their surface, is very similar to that of bioglass, their interactions with the human body should be comparable. Like bioglasses, silica-based mesoporous materials should develop a nanostructured apatite layer on the material surface when it comes in contact with physiological fluids,

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Figure 4.7: Molecular sizes of common medications that may be good candidates for drug delivery (republished from reference [24] with permission from Wiley, Inc. and conveyed through Copyright Clearance Center, Inc.).

and the apatite layer should be comparable to biological apatite and support bonding between implants and bone tissue [2]. The formation of such an apatite layer is also related to silanol groups (Si–OH), which act as nucleation sites. In addition to silanol groups, porosity influences apatite layer formation. The nucleation rate is affected by both pore size and connectivity; that is, the larger and more accessible the pores, the higher the ion diffusion and (consequently) the faster the apatite layer formation. Therefore, mesoporous silica is advantageous because of its adjustable pore diameters and volumes, high surface area, high silanol concentration, and possible bioactivity owing to these properties [22]. Silanol groups are formed because of the lack of bonding partners for Si–O− groups, which are the result of SiO44− tetrahedrons bonded by shared oxygen atoms to fewer than four surrounding tetrahedrons. Therefore, the concentration of these connectivity defects determines the silanol group concentration. Owing to structurally disarrayed tetrahedral building blocks, the number of connectivity defects is very high. Therefore, this parameter plays a major role in the material behaviour because the material hydrophilicity increases with increasing silanol group concentration. Additionally, silanol moieties can be functionalised with organic groups (e.g. Si–O–R) to enable interactions with other compounds such as drugs. The number of connectivity-defect-induced silanol groups is especially influenced by the surfactant removal. For example, heat treatment (calcination) condenses silanol groups, which decreases the number of defects. In contrast, surfactant solvent extraction does not affect the defect concentration [24].

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The initial practical tests of silica-based mesoporous material bioactivity indicated that parameters such as pore size, structure, texture, and material composition all influenced the apatite formation rate, as determined by studying the bioactivity of SBA-15, MCM48, and MCM-41 (SBA = Santa Barbara Amorphous; MCM = Mobil Composition of Matter) mesoporous silica materials, which exhibit unique pore diameters and volumes, silanol concentrations, and surface areas (Table 4.4) in SBF solutions. These silicas were dipped in SBFs to investigate apatite layer formation. An HCAp layer formed on the MCM-48 surface after 60 d, SBA-15 required only half that time, and no apatite layer formed on the MCM-41 surface (Figure 4.8) [22]. To determine the different biological behaviours of these silicas, their respective properties must be examined in more detail. SBA-15, which exhibited the largest pore diameter, also exhibited the fastest HCAp formation compared to MCM-48 and MCM-41. SBA-15 also exhibited additional micropores (e.g. porosity = 0.061 cm3/g), which positively affected apatite layer formation. Therefore, MCM-48 and MCM-41 should exhibit similar behaviours because they exhibit the same pore diameter of 3.6 nm. However, the large difference in bioactivity after 60 d (especially no bioactivity for MCM-41 in this time range) could be explained by their unequal silanol concentrations. As previously mentioned, silanol moieties can act as nucleation sites; therefore, the much higher MCM-48 silanol concentration compared to six times lower concentration for MCM-41 explains their different bioactivities for the same pore size. Furthermore, the MCM-41 pore volume is slightly lower than the MCM-48 one, which probably negatively affects bioactivity [28]. Therefore, pore volume and reachability both play a role in ion diffusion, which is higher for larger and well-connected pores and leads to a higher nucleation rate. In addition, the lower MCM-41 bioactivity can be explained by the two-dimensional hexagonal structure, whereas MCM-48 exhibits a three-dimensional cubic scaffold and channel connections, thereby resulting in higher ion diffusion [22]. Table 4.4: Physical properties of SBA-15, MCM-48, and MCM-41 [22, 28]. Material Surface area (m/g) Pore size Pore volume Silanol concentration (− mmol/m) (nm) (cm/g) SBA- MCM- MCM-

  

. . .

. . .

. . .

Introducing a porous structure to HAp theoretically overcomes some of the technical difficulties of non-porous HAp; that is, relatively low bonding surface areas, high material density [29], and slow replacement of implant structures with natural bone tissue [12]. Thus, implants can act as templates that can be filled with new bone tissue. Pore size influences bone regeneration by offering space for cells to attach and bone tissue to grow into the implant. Another important factor is pore interconnectivity because more interconnected pores ensure cellular distribution and vasculo-genesis and

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Figure 4.8: Influence of SBFs on SBA-15, MCM-48, and MCM-41 surfaces over two months. For each compound, left and right images are SEM micrographs and corresponding EDS spectra, respectively (republished from reference [22] with permission from Royal Society of Chemistry and conveyed through Copyright Clearance Center, Inc.).

improve bone in-growth. For good bone in-growth and vascularisation, pores must be at least 100 µm in diameter. However, mechanical strength decreases with increasing porosity [29]. Therefore, conventional porous HAp (which exhibits low pore interconnectivity and non-uniform pore geometry) hinders bone tissue in-growth, whereas HAp with highly interconnected pores enables osteoconduction into the centre of implants. For complete tissue regeneration, pore size, density, interconnections, and shapes are crucial factors [12].

4.3 Synthesis methods for nanostructured bioceramics Methods used to synthesise nanostructured bioceramics vary depending on the ceramic composition, desired crystal morphology (e.g. needles, rods, spheres, plates, flowers, bowknots, fibres, strap-like, or mesoporous structures (Figure 4.9)), technological

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Figure 4.9: Different morphologies of calcium-phosphate-based nanoparticles synthesised using (A, B) sol–gel processing, (C) co-precipitation, (D) emulsion technique, (E) hydrothermal processing, (F) ultrasonication, (G) mechanochemical method, (H–L) templating, (M) microwave processing, (N) combined emulsion and hydrothermal synthesis, and (O) combined microwave and hydrothermal synthesis (reprinted from reference [30] with permission from Elsevier).

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difficulty, equipment requirements, reaction time and conditions, cost, and product yield [30]. Bioceramic syntheses like that of calcium phosphates can be categorised as solidand wet-state methods. Solid-state reactions involve heating pure reactant powders to high temperatures (usually >700 °C) for long periods of time, are expensive to conduct, and the products they generate tend to be highly stoichiometric and exhibit wellcrystallised structures. Furthermore, because HAp particles sinter at high temperatures, the product mostly consists of sintered bulk. Thus, the product must be crushed into smaller particles (usually by ball milling) to obtain nanocrystals [11]. Wet-state methods are conducted in solvent solutions at relatively low temperatures. Products obtained from wet-state reactions are frequently naturally nanoparticles but exhibit poor stoichiometry and crystallinity [11]. The different methods are distinguished by the required effort, sample treatment, reaction times, conditions, yield, etc. The sol–gel process uses simple devices to provide high-purity homogenous crystals but requires a considerable amount of time and high sintering temperatures. In contrast, co-precipitation is low-cost and although yields crystals exhibiting bone-like structures, the crystals only exhibit poor crystallinity. Therefore, method advantages and disadvantages should be considered before preparing the required product. Despite ultrasonication usually produces low yields and requires special devices, it provides low particle aggregation, narrow diameter distribution, and fast crystallisation at low temperatures. Even when reaction conditions are easily achievable using simple devices and the product exhibits excellent properties, product yield determines whether the synthesis method is worthwhile [30]. To emphasise the complexity of selecting the most appropriate synthesis method to obtain the desired product, Table 4.5 compares and contrasts the characteristics of various synthesis methods available for producing calcium phosphate-based nanobioceramics; namely, sol-gel processing, co-precipitation, emulsion technique, hydrothermal processing, mechanochemical method, ultrasonication, templating, microwave processing, combined emulsion–hydrothermal synthesis, and combined microwave–hydrothermal synthesis [30]. Based on the advantages and disadvantages of the available synthesis methods, coprecipitation and sol–gel processing were considered particularly well-suited for producing medical-grade calcium-phosphate-based nanocrystals and will be explained in detail.

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Table 4.5: Synthesis methods, advantages, disadvantages, and morphologies of obtained calciumphosphate-based nanoparticles (reprinted from reference [30] with permission from Elsevier). Synthesis method

Nanocrystal morphology

Advantages

Disadvantages

Sol-gel processing

–

– –

High purity Homogenous components Simple devices required

–

Crystals exhibiting bone-like structures and components Low cost Particle aggregation can be controlled using biomacromolecules

– –

Low crystallinity Particles aggregate easily

Low aggregation Narrow diameter distribution Particle size and morphology can be easily adjusted

–

High sintering temperatures Time consuming Complex reaction system

Aggregated

– Co-precipitation

– – –

Needle-like Spherical Rod-like

–

– –

–

High sintering temperatures Time consuming

– – – –

Needle-like Spherical Nanorods Nanoplates

– –

– – – –

Needle-like Spherical Fibrous Nanorods

– – – –

High crystallinity High purity Good dispersibility Good control of particle size and morphology

–

Mechanochemical method

– –

Fibrous Spherical

– – –

Simplicity Low cost No high-temperature sintering required

– –

Particle aggregation Special devices are required

Ultrasonication

– –

Needle-like Spherical

–

Low particle aggregation Narrow diameter distribution Rapid crystallisation Low temperatures

–

Special devices are required Low yield

Emulsion technique

Hydrothermal processing

–

– – –

– –

– –

–

Special devices are required Time consuming Very low yield

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Table 4.5 (continued) Synthesis method

Nanocrystal morphology

Advantages

Disadvantages

Templating

– – –

Nanorods Nanoplates Mesoporous and fibrous nanostructures Flower and layered structures

–

–

Bowknot Rod Needle-like Flower structures

– –

– Microwave processing

– – – –

– – – –

Spherical Rod Fibrous Strap-like structures

High crystallinity Narrow diameter distribution Good control of particle size and morphology Rapid synthesis

–

–

–

Low particle aggregation Narrow diameter distribution Good control of particle size and shape Low temperatures

– – –

High crystallinity Rapid synthesis Low temperatures

–

–

– – –

Microwave– hydrothermal combination

– –

Needle-like Spherical

–

– –

– Emulsion– hydrothermal combination

Good control of particle morphology, size, structure, tropism, and array

–

– – –

–

Requires templating agents High sintering temperatures required Time consuming Low yield Special devices are required No large-scale synthesis

Special devices are required Time consuming Complex reaction Very low yield

Special devices are required Very low yield

Co-precipitation This wet-state method is cost-effective [30], easy to perform, and requires mild reaction conditions. Typically, the co-precipitation synthesis of calcium-phosphate-based nanobioceramics (like HAp) begins with a mixture of two simple aqueous solutions, one containing calcium ions; the other, ortho-phosphate ions. Either each solution has a pH > 7 or the pH was increased for the mixture to produce a supersaturated HAp solution and coprecipitate HAp nanoparticles using calcium hydroxide and phosphoric acid (eq. (4.1)) or calcium nitrate, diammonium hydrogen phosphate, and ammonium hydroxide (eq. (4.2)), as follows [11]: 10CaðOHÞ2 + 6H3 PO4 ! Ca10 ðPO4 Þ6 ðOHÞ2 # + 18 H2 O

(4:1)

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10CaðNO3 Þ2 + 6ðNH4 Þ2 HPO4 + 8NH4 OH ! Ca10 ðPO4 Þ6 ðOHÞ2 # + 20NH4 NO3 + 6H2 O (4:2) Although this method is highly efficient, it is not without disadvantages. HAp crystals obtained by co-precipitation are prone to aggregation, which can be counteracted by adding polymers and surfactants as dispersants to the reaction mixture. However, these organic additives then become integrated into the compound structure. Applying high sintering temperature improves crystallinity and removes remaining additives. [30]. Moreover, the influences of synthesis conditions such as temperature, reaction time, calcium ion concentration, calcination, and the reactants employed during coprecipitation have been extensively discussed in the literature [4]. Special attention has been dedicated to study the parameters affecting apatite nanoparticle morphology, which depend on reactant concentration, ionic strength, temperature and solution pH [11]. Strictly controlling synthesis parameters is important for achieving the desired particle size, which reportedly increases linearly with increasing reaction

Figure 4.10: SEM images of calcium-phosphate-based nanoparticles co-precipitated at initial pH values (a–c) 10 and (d–f) 6 and at (a, d) 25 °C, (b, e) 50 °C, and (c, f) 80 °C (reprinted from reference [31]).

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temperature, and for guaranteeing high dispersibility by applying an appropriate solvent, dispersant, and product drying method [4]. Figure 4.10 shows the morphology of calcium phosphate particles co-precipitated at different pHs and temperatures. For example, the morphology of HAp particles coprecipitated under alkaline conditions changed from spherical (17 nm; Figure 4.10a) to rod-shaped (154 × 13 nm; Figure 4.10b) to elongated rod-shaped (585 × 43 nm; Figure 4.10c) with increasing reaction temperature. Under initially acidic conditions, OCP and DCPD microparticles co-precipitated exhibiting a blade-shaped morphology at lower temperatures (Figure 4.10d and e, respectively), whereas HAp particles coprecipitated at a higher temperature and exhibited an elongated fibre morphology (3,550 × 48 nm; Figure 4.10f) [31].

Sol-gel processing A gel is a compound exhibiting low mechanical strength and is composed of at least an extended solid network filled with either a liquid or a gaseous phase. The network filler determines the type of gel. For example, hydrogels contain water as a liquid phase, and alcogels and xerogels are filled with an alcohol and a gaseous phase, respectively [32]. Colloidal solutions are sols wherein particles are dispersed in liquids so that particle interactions are reduced. However, small particles aggregate into larger ones and decrease the particle surface area-to-volume ratio. Therefore, to prevent particle aggregation and subsequent precipitation, colloidal solutions must be stabilised by changing particle surface charges, which can be adjusted by increasing or decreasing pH. Hence, sols are gel precursors. In particle-charge-stabilised sols, changing the pH in the opposite direction can collapse the sol, and particle agglomeration (e.g. condensation between particles and hydrogen bonds) forms an irregular gel network [32]. Calcium and phosphorus alkoxides can be used to synthesise calcium phosphates. The solid product obtained using sol–gel synthesis then contains ACP intermediates and unreacted precursors and is subsequently sintered at 400–500 °C, which is sufficient to produce well-crystallised HAp because HAp powder is usually sintered in the range 800–1,000 °C. As the reaction produces a wide-network gel that is subsequently sintered, the obtained product is polycrystalline. To produce nanoparticles, the raw product must be often ground or milled [11].

4.4 Functionalisation of bioceramics In general, one important objective of functionalising bioceramics is to improve their biocompatibility and bioactivity. In this context, porous materials offer unique

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functionalisation opportunities due to their high inner surface areas. For this reason, the functionalisation of mesoporous silica is discussed in this section in details. For instance, the bioactivity of the mesoporous silica MCM-41 is reported to be improved by either doping the scaffold with phosphorous or functionalising it with bioglass. This way, phosphorus-containing MCM-41 exhibits an ordered mesoporous scaffold in which Si–O–P bonds are in the outer particle parts rather than at the inner core. Hence, MCM-41 textural and mesostructural properties are preserved even at low phosphorus contents of approximately 1%. When exposed to simulated body fluids, MCM-41 develops an apatite-like layer after two weeks. Using bioglass (SiO2–CaO– P2O5) instead of phosphorous, the in vitro bioactive response of the biphasic mixture is that an apatite layer forms after only few hours and finishes forming after one day. The bioglass acts as a nucleation site to accelerate apatite layer formation and preserve both textural and mesostructural properties [22]. In addition, bioceramic surfaces are suitable for not only adsorbing advantageous cells and proteins but also providing a bacterial habitat. Surfaces that prevent random protein adsorption and the formation of bacteria-loaded biofilms are called “anti-fouling surfaces”. Such surfaces can prevent implant infections, which would hinder healing. Anti-fouling surfaces can be produced by attaching zwitterionic polymers to surfaces [2]. As zwitterionic polymer moieties are charged while the overall polymer is electrically neutral, a hydration layer is formed through electrostatic interactions and hydrogen bonding and strongly repels proteins that approach the surface [32]. To achieve this anti-fouling property, mesoporous glasses could be modified with carboxyl and amino groups. Figure 4.11 shows one example in which bacteria colonised the unmodified glass surface (SBA-15), and even more bacteria colonised the carboxyl-modified glass (SBA15CES). Amino-group modification (SBA15NH2), on the other hand, slightly decreased bacterial adhesion, and a mixture of carboxyl and amino groups (SBA15APTES/CES) prevented nearly all bacterial colonisation (Figure 4.12) [2]. Mesoporous silica surfaces can be additionally functionalised to control drugrelease rates. For example, silanol groups can replace H atoms with organic radical groups (R) to functionalise silica scaffolds and can be further modified if R contains reactive groups to generate a large family of hybrid materials. In silanol functionalisation, the silanol group and R-containing precursor condense to form a new functional group according to the following reaction [24]: Si−OH + X3 Si−R−Y ! Si−O−SiX3−n −R−Y + nHX.

(4:3)

The chemical species usually used in this reaction are X = Cl, OMe, OEt; R = alkyl chain; and Y = OH, SH, NH2, SO3H, Cl, F, CH3, Ph, etc. [24] This process can be applied post-synthetic to produce highly functionalised pore walls. After porous silica is synthesised, the functionalising precursor is added and reacts with pore wall silanol groups. Silica scaffolds can also be functionalised using a one-pot method in which a silica scaffold precursor and silicon alkoxy functional groups react simultaneously, wherein the alkoxy groups are hydrolysed and condensed with the silica precursor.

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Figure 4.11: SEM images of mesoporous silica surfaces (SBA-15) exhibiting different functional-groupdependent anti-fouling properties. Rod-shaped bodies surrounded by fibres are E. coli bacteria attached to surfaces (republished from reference [2] with permission from Wiley, Inc. and conveyed through Copyright Clearance Center, Inc.).

Figure 4.12: Concentrations of adsorbed proteins and corresponding colony-forming units of bacteria attached to different functionalised surfaces (republished from reference [2] with permission from Wiley, Inc. and conveyed through Copyright Clearance Center, Inc.).

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After the reaction, the functional groups are located outside and inside the walls, resulting in less functionalisation than post-synthetic functionalisation [24]. Consequently, the mesoporous structural order decreased with increasing organo-silica precursor content in one-pot synthesis. Therefore, it is important to notice that the organic functional group content must be ≤ 40 mol% to preserve the mesoporous structural order [26]. An additional practical example of how functionalised silica scaffolds influence drug delivery rates is shown by aminopropyl-moiety-functionalised MCM-41 scaffolds [24]. When the scaffolds were loaded with ibuprofen, the uncontrolled drug-release rate decreased by approximately five times compared to the drug-release rate of unmodified MCM-41 scaffolds because the interaction between the scaffold amino and ibuprofen acid groups decreased drug mobility. Hence, the scaffold functionalised with more polar groups exhibited higher ibuprofen adsorption than the scaffold exhibiting non-polar groups. In contrast, ibuprofen exhibits very high mobility in unmodified materials. Therefore, functional groups can be chosen to control delivery rates for specific drugs [24]. Moreover, non-polar moiety functionalisation can not only improve scaffold–drug interactions but also decrease wettability, which hinders water molecules from penetrating pores and reduces drug release [26]. On the other hand, the use of calcium phosphates in load-bearing applications is limited because mechanical properties of calcium phosphate deteriorate with increasing porosity [29]. However, by functionalising this calcium phosphate with zirconia (ZrO2), the bending and compressive strength, fracture toughness, and elasticity modulus can be improved. Zirconia-toughened via coating HAp particles exhibits mechanical strength equal or superior to that of human compact bone [33]. In vitro and in vivo tests both showed that although these compounds are biocompatible, dissolution is slower than conventional HAp in SBFs [29]. Similar to the functionalisation of silica surfaces to obtain antibacterial properties [2], HAp can be modified to acquire these properties and can be implemented by adding zinc oxide (ZnO) to the HAp scaffold. Here, zinc is responsible for the antibacterial properties and is also important for controlling cell proliferation. In addition, zinc may be relevant for accelerating bone repair, thereby improving the scaffold in different ways. Such a scaffold can be obtained by soaking porous HAp in a solution containing dispersed ZnO particles, drying the compounds, and then slowly heating the material to the sintering temperature. Reportedly, the obtained material contains zinc-substituted HAp in which the substitution yields a non-stoichiometric product according to the formula Cax Znx Ca5 ðPO4 Þ3 ðOHÞ and exhibits good biocompatibility and solubility in bodily fluids, while preserving three-dimensional porosity [34].

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4.5 Applications of bioceramics Current research is trending away from bioinert bioceramics towards bioactive ones. As bioinert bioceramics do not bond directly to bone or stimulate bone growth, many implants produced using these materials persist for only approximately 15 years. Using bioactive materials can extend implant lifetimes and enable implants to be better accepted by the human body, while interacting with the organism [30] and opening the doors to various applications. One of the most intuitive bioceramic applications is bone tissue engineering. These materials can replace damaged bone to temporarily preserve mechanical stability while new bone tissue is being formed. For such applications, these materials must fulfil several requirements to be appropriate for implant production. For example, for materials to be used as a template for new bone tissue growth, they must exhibit mechanical properties comparable to those of the replaced tissue, and the materials must be resorbed by the host tissue at a similar rate as the new bone is forming while supporting bonding with the host bone, not causing scar tissue, and being biocompatible. Finally, they must be processable into any shape to match bone defects, promote cellular adhesion and activity, and match international standard clinical application requirements [2]. In addition to bioceramics, metal implants produced using stainless steel and titanium and its alloys are usually used to replace damaged or lost bone. Due to their inertness, corrosion resistance, load-bearing mechanical properties, and biomaterial compatibility, titanium-based compounds have been extensively applied to bone replacements, which opens the door to combining titanium-based alloys and bioceramics such as HAp-coated titanium implants. Biocompatible and bioactive coatings (commonly applied using plasma or arc plasma spraying) improve metal implants. For example, biocompatible and bioactive coatings protect implants from corrosion and contact with bodily fluids and support surface cellular adhesion, prosthetic fixation, and damaged tissue healing. However, because such coatings are brittle, the coating thickness cannot be arbitrary and is limited to < 70 µm [35]. In tissue engineering, macroporous HAp foams synthesised using sol–gel processing reportedly have been applied to bone-growth scaffolds exhibiting porosities of 10–15 nm and 1–400 µm, which is very similar to the hierarchical porosity of meso- and macropores in natural bone tissue. Dip-coating these foams with biopolymers such as glutaraldehydecross-linked gelatine and polycaprolactone improved the foam mechanical properties (e.g. flexibility and resistance) through water contact and foam swelling while retaining the original foam porosity. Notably, in bioceramic medical applications, heavy metal ions can be captured by macroporous HAp foams, which are suitable for treating ingestion-induced heavymetal intoxication through intestinal tract absorption. Heavy metals are absorbed by ionically substituting Ca2+ ions in the HAp scaffold, thereby trapping heavy metal ions such as Ba2+, Sr2+, Cd2+, Pb2+, Zn2+, and Cu2+ within the foam. Under very acidic

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conditions, such as those in the stomach (pH 1.2), the HAp must be coated with biopolymers because of its high solubility under these conditions. HAp modified like this scaffolds rapidly and effectively captured heavy metal ions and did not release them under basic conditions. Metal ion capture depends on the biopolymer hydrophilicity. In in vitro tests, captured Pb2+, Cd2+, and Cu2+ concentrations increased to 405, 316, and 378 µmol/g, respectively, suggesting that these low-cost, easily handled bioceramic materials could also be used to purify heavy-metal-contaminated water [2]. Bioceramics have also been prepared by combining ferromagnetic bioactive glass with a magnetic glass ceramic for application to hyperthermia-induced cancer treatment [17]. In an external alternating magnetic field, such materials absorb magnetic energy and release it as heat. The required temperature range is 43–47 °C, at which cancer cells are almost irreversibly and selectively destroyed. Moreover, such materials could replace cancerous bone tissue and prevent any remaining cancerous cells from metastasising. Additionally, bioceramics can be loaded with drugs to further improve cancer treatment. Hence, this effective localised treatment is advantageous compared to cancer treatment with chemotherapy, which is usually accompanied by low tolerance levels, toxicity, and low efficacy [2]. One important drawback of this approach is that, at hyperthermia temperatures, not only cancerous tissue can be targeted and therefore also overheating of the surrounding healthy tissue may occur. Although hyperthermia is not an ideal stand-alone cancer treatment, it has shown high efficiency in combination with conventional radiotherapy and chemotherapy. In an experiment to treat murine mammary breast carcinoma, a lithium-ferrite-containing haematite (Fe2O3) matrix embedded in glassy SiO2– P2O5 was implanted, and the implant reduced tumour re-growth by ~ 50 % [2]. Because bioceramics can be loaded with pharmaceutical compounds, they can be used to release drugs inside the human body. Potential materials for application to drug-delivery systems include calcium carbonate [6], mesoporous silica [26], and calcium phosphate wherein drugs can be directly incorporated into particles, adsorbed on particle surfaces, or both. Particle surface area plays a major role in drug adsorption, which favours porous materials exhibiting very high surface areas per gram [36]. The slow degradation of some bioceramics, on the other hand, is beneficial for continuously releasing drugs over certain extended periods [6]. A common problem in systemically treating diseases or infections is the frequent requirement of high drug doses to ensure that the necessary dose reaches the target. As such treatments can cause harmful side effects owing to the overdose risk, drugs must be selectively locally released [6, 36]. Such drugs can support, for instance, osteogenic healing, prevent infections with the help of antibiotics, or directly treat tissueimpairing diseases. Hence, low bone tissue blood flow makes localised treatment more advantageous than systemic treatment because effective drug concentrations cannot otherwise be reached. Additionally, fewer drugs are required at lower concentrations to ensure effective drug doses reach the target area [36].

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To prevent bone tissue infections, bioceramics can be also loaded with and continuously release antibiotics. Therefore, bacterial adhesion can be reduced at implantation sites, thereby decreasing the risk of infection. Bioceramics can also be loaded with drugs to treat diseases such as osteoporosis, wherein bone density and mass both decrease causing bone fractures or osteosarcoma, which is a type of bone tumour. Furthermore, bone cancer patients often suffer from high chemotherapy-induced toxicity, tissue removal, or additional radiotherapy-induced bodily stress. Therefore, localised cancer treatment is more beneficial because anti-cancer agents such as cisplatin can be easily adsorbed by HAp nanoparticles and administered at low concentrations, because they can be very toxic at high doses [36].

4.6 Conclusions and outlook Since the beginning of their development, bioceramics have massively evolved in the ensuing decades starting with first-generation bioceramics, which were bioinert and were used to replace damaged tissue, thereby preventing any interaction with surrounding tissues. In contrast, second-generation bioceramic materials were developed to exhibit bioactive behaviours to repair damaged tissue [2]. With the development of third-generation bioceramics, materials support bodily self-repair [3]. Therefore, because bioceramics are biocompatible, they are extensively used in medical applications [2]. Bioceramic properties can be improved by scaling micro- to nanodimensions [13] especially because bioceramic nanoparticles exhibit higher surface area-to-volume ratios and consequently enhanced reactivity compared to their bioceramic microparticle counterparts [4]. Porous materials such as bioglasses [26], HAp [2] and silicas [28] exhibit high surface areas [24], which favours bone in-growth [12]. Bioceramics can be prepared using different chemical compositions such as calcium phosphates (e.g. bone-like HAp) [12], calcium carbonates [6], and silica-based materials [28]. Calcium phosphates exhibit good biocompatibility and bioactivity especially because of their similarity to the bone inorganic phase. However, calcium phosphate is not restricted to hard-tissue-based medical applications and can be applied to soft tissues such as blood vessels, nerve tissues, and cartilage [5]. In addition to being low-cost and readily abundant, calcium carbonates exhibit good bioresorptivity in which pH sensitivity and slow degradation under physiological conditions open the doors to possible drug-delivery applications [6]. In bioglasses, additives such as Na2O, CaO, and P2O5 can induce the formation of an apatite-like surface layer, which can bond to bone structures [18]. Thus, BG implants are resorbed and replaced by new bone tissue [2].

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Several synthesis methods ranging from co-precipitation to microwave processing have been used to prepare nanostructured bioceramics, depending mainly on the required crystallinity or morphology [30]. Synthesised bioceramics can be functionalised to improve their properties by adding dopants to form hydroxycarbonate apatite layers on mesoporous silica surfaces [22], to improve mechanical properties with coatings [33], or decreasing bacterial adhesion on functionalised bioceramic surfaces [32]. As bioceramics positively interact with the human body, most research has focused on medical applications [36]. For instance, bioceramics act as scaffolds for regenerating bone in tissue engineering [2] or as coatings to increase bodily acceptance of metallic implants [35]. In cancer treatment, hyperthermia-based therapy can be performed using ferromagnetic bioactive glasses [17]. Bioceramics are also suitable for developing drug-delivery systems in which drugs can be trapped inside bioceramic particles [6] or adsorbed on bioceramic particle surfaces [26]. Inside the body, drug-containing bioceramic-based implants can constantly deliver drug loads and support bone regeneration [36] or cancer treatment [2]. Currently, although research on bioceramics mainly focuses on bone and tooth applications, additional investigations are required for bioceramics to interact with soft tissues and for healing wounds. Another important research topic is elucidating the mechanisms through which bioceramics interact with body tissues. Understanding how different cells respond to bioceramic compositions and how multiple cell generations synchronise with the stimuli of bioceramic materials would greatly expand the application potential of these materials [37].

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[29] R. Shao, R. Quan, L. Zhang, X. Wei, D. Yang, S. Xie, Porous hydroxyapatite bioceramics in bone tissue engineering: Current uses and perspectives. J. Ceram. Soc. Jpn. 123, 17–20, 2015, (and references therein). [30] Y. Hong, H. Fan, B. Li, B. Guo, M. Liu, X. Zhang, Fabrication, biological effects, and medical applications of calcium phosphate nanoceramics. Mat. Sci. Eng. R. 70, 225–242, 2010, (and references therein). [31] M. Okada, D. Hiramatsu, T. Okihara, T. Matsumoto, Adsorption and desorption behaviors of cetylpyridinium chloride on hydroxyapatite nanoparticles with different morphologies. Dent. Mater. J. 35, 651–658, 2016, (and references therein). [32] a) M. Colilla, I. Izquierdo-Barba, S. Sánchez-Salcedo, J. L. G. Fierro, J. L. Hueso, M. Vallet-Regí, Synthesis and characterization of Zwitterionic SBA-15 nanostructured materials. Chem. Mater. 22, 6459–6466, 2010, (and references therein). b) H.K. Schmidt, Das Sol-Gel-Verfahren: Anorganische Synthesemethoden. Chem. unserer Zeit 35, 176–184, 2001. [33] T. Matsuno, K. Watanabe, K. Ono, M. Koishi, Microstructure and mechanical properties of sintered body of zirconia coated hydroxyapatite particles. J. Mater. Sci. Lett. 19, 573–576, 2000, (and references therein). [34] C. A. Martínez, U. Gilabert, L. Garrido, M. Rosenbusch, A. Ozols, Functionalization of hydroxyapatite scaffolds with ZnO. Procedia Mater. Sci. 9, 484–490, 2015, (and references therein). [35] M. Prakasam, J. Locs, K. Salma-Ancane, D. Loca, A. Largeteau, L. Berzina-Cimdina, Fabrication, properties and applications of dense hydroxyapatite: A review. J. Funct. Biomater. 6, 1099–1140, 2015, (and references therein). [36] D. Arcos, M. Vallet-Regí, Bioceramics for drug delivery. Acta Mater. 61, 890–911, 2013. (and references therein). [37] L. L. Hench, The future of bioactive ceramics. J. Mater. Sci. Mater. Med. 26, 86, 2015, (and references therein).

Huayna Terraschke✶, Markus Suta

Chapter 5 In situ luminescence analysis of coordination sensors (ILACS): looking inside chemical reactions Monitoring the formation of solid materials independent of their crystallinity and of synchrotron radiation at common university laboratories is a challenge. A solution to this issue can be offered by the new in situ luminescence analysis of coordination sensors. With this approach, lanthanide ions are incorporated as local sensors to the structure of the monitored compounds. As the optical properties of lanthanide ions are strongly influenced by the coordination environment, the structural evolution of the monitored compounds can be tracked during synthesis by measuring real-time luminescence spectra under real reaction conditions.

5.1 Introduction The synthesis of new functional materials with desired physical properties or a welldefined chemical reactivity is still a major field in solid-state chemistry. Many technologies largely depend on the development of new materials with improved properties. In this context, solid-state compounds represent a significant part of it and, therefore, form the basis for many industrial areas. Indeed, there are chemical rules and laws at hand but the outcome of a reaction cannot be predicted with certainty. Typically, reactants are mixed and reacted using different synthesis methods, the products are isolated, and routinely characterised. In most cases, very little is known about the mechanisms of such reactions beginning with the nucleation up to the formation of long-range crystalline structures. Due to this lack of knowledge the development of rational synthetic strategies or the discovery of new compounds is difficult to achieve. However, this problem might be solved by analysing not only the readily prepared product but also investigating the different stages of product formation, for example, nucleation and crystal growth in view of structural changes or the formation of possible reaction intermediates that might be of interest. Such intermediates can be stable or metastable and they can transform into each other very fast, and therefore, can easily be overlooked using conventional ex situ techniques. In this context, in situ monitoring of chemical reactions has become of increasing interest [1–3].

✶

Corresponding author: Huayna Terraschke, E-Mail: [email protected]

https://doi.org/10.1515/9783110459098-005

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For the in situ monitoring of chemical reactions or phase transitions, several characterisation techniques have been applied, which include, for example, X-ray diffraction (XRD) [4–8], pair distribution function (PDF) analyses [9, 10], X-ray absorption spectroscopy (XAS) [11, 12], Raman spectroscopy [13], infrared (IR) spectroscopy [14], nuclear magnetic resonance (NMR) spectroscopy [15], and mass spectrometry (MS) [16]. Although powerful methods, they cannot be applied to any problem. XRD-based methods, for example, are not appropriate for analysing very light atoms with low electron density, amorphous materials, or crystalline phases in low concentration. While in situ XAS delivers information about the short-range order and electronic state of the absorber, synchrotron radiation is required, limiting accessibility of such experiments. IR and Raman spectroscopy allow monitoring changes in chemical bonding situations, but are more suitable for organic compounds. NMR, on the other hand, probes the local environment of NMR active atoms, thus being restricted to such atoms and being more appropriate for reactions in liquids, which do not require a high time resolution. In situ luminescence analysis of coordination sensors (ILACS) is a new analysis technique developed at the Kiel University to monitor and elucidate the mechanisms of chemical reactions. Within this technique, lanthanide (Ln) ions [17–19] are employed as local coordination sensors, because of their spectral sensitivity to their coordination environment. Changes in the surroundings of the coordination sensors are detected by submerging an optical fibre into the reactor content, transporting the emitted light to a fast charge-coupled device (CCD)-based detector and recording luminescence spectra with very high time resolution. Previously reported studies on time-dependent or in situ luminescence measurements are mostly limited to luminescence titration, temperature-dependent luminescence measurements, or quantification of luminescence lifetime. Luminescence titration is generally used to quantify the compounds present in a solution with unknown concentration by detecting the reactions of these compounds with emitting markers, generating a concentration-dependent luminescence signal. Temperature-dependent measurements are typically used to investigate quenching effects during heating processes or during the improvement of the spectral resolution by cooling and decreasing the energy loss caused by thermally induced non-radiative electronic transitions. Time-resolved luminescence measurements are used to determine the decay kinetics from the excited states. Nevertheless, these techniques provide no mechanistic information about the formation of solid materials. Most in situ analysis of optical properties previously available for tracking chemical reactions were primarily based on the time-dependent measurements of UV/Vis absorption properties, which afford valuable information on the formation of solid materials, such as changes in the bandgap energy. However, local changes in the coordination environments of ions are not accessed by these techniques, and they do not provide information about the short-range atomic arrangements of the formed compounds. To date, information about the changes in the first coordination spheres of lanthanide and actinide ions is also acquired via time-resolved laser-induced fluorescence spectroscopy (TRLFS) [20, 21].

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Section 5.2 explains the working principle of ILACS in detail, also differentiating it from TRLFS. In Section 5.3, the theory behind the optical properties of lanthanide coordination sensors is shortly reviewed. Section 5.4 describes the state-of-the art instrumentation available for conducting ILACS experiments, while examples of the study cases using these instruments are mentioned in Section 5.5. Subsequently in Section 5.6, applications of the ILACS experimental instrumentation for monitoring chemical reactions without coordination sensors are discussed. Finally, Section 5.7 summarises the chapter and offers an outlook on the future of this technique. For additional information on this subject, our previous review article can be referred [3].

5.2 Working principle Basic ILACS involves the introduction of a suitable local coordination sensor into the structure of the material to be investigated. This can be performed by either producing an Ln-based compound itself or incorporating extremely small amounts of Ln ions into the studied host compound of interest (Figure 5.1). Coordination-sensitive spectroscopic changes of the coordination sensors are detected during the reaction. For this purpose, an optical fibre is submerged into the reaction vessel, transporting the emitted light to a fast CCD-based detector. This approach is advantageous for complementing techniques such as XRD, especially because it is able to characterise not only highly crystalline compounds, but also amorphous materials and small crystallites. In addition, ILACS is able to characterise ions in solution and monitor desolvation processes. Compared to general synchrotron radiation-dependent in situ methods, ILACS can be also usefully complementary, because it can be conducted in conventional university laboratories, increasing its availability (Figure 5.2). The high frame rate, which reaches several spectra per second depending on the measured emission intensity of the synthesised compound and required integration time, makes ILACS significantly advantageous in comparison to, for instance, in situ NMR spectroscopy. High time resolution is important for examining fast processes such as nucleation and detection of short-lived intermediates. Intermediates may present interesting properties for the development of new materials and can remain undetected when only ex situ methods or methods with low time resolutions are applied. As comprehensively explained in Section 5.4, ILACS is performed in an external reactor system, which enables the reproduction of real reaction conditions, including stirring, heating, cooling, dosing, as well as the potential application of ultrasound or generation of high pressure within the reaction vessel. External reactor systems also allow the modification of used setups for both stationary and mobile use. Mobile experimental setups enable multimodality experiments, combining ILACS with synchrotron-based in situ X-ray powder diffraction, besides with measurements of pH, ion conductivity, redox potential, or IR spectroscopy for confirming and complementing the in situ luminescence results.

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Figure 5.1: Working principle of ILACS. Lanthanide ions are incorporated into the material to be analysed, providing an emission spectrum. Changes in the crystal structure during reaction result on the modification of the optical properties, which can be monitored by in situ luminescence measurements (reproduced from reference [49] with permission from the Chinese Chemical Society (CCS), Peking University (PKU), and the Royal Society of Chemistry).

This is an important difference between ILACS and TRLFS, which is generally conducted in cuvettes under stationary conditions. In addition, ILACS uses mostly a combination of an Xe lamp and a monochromator instead of a laser as the excitation source, which thereby increases the flexibility of the applied excitation wavelength and facilitates the investigation of a larger range of electronic transitions and selective excitation of different doping sites.

Figure 5.2: Overview of the features offered by ILACS.

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5.3 Spectroscopic background of coordination sensors This section outlines the spectroscopic foundations of the electronic transitions used by the lanthanide coordination sensors as well as the types of structural information that can be obtained by critically examining the spectral emissions of these ions. Luminescence resulting from 4fn ↔ 4fn−15d1 or 4fn ↔ 4fn electronic transitions is caused by the partially filled 4f orbitals of elements ranging from Ce to Yb. Classical examples of these electronic transitions are observed for divalent and trivalent Eu ions. Neutral Eu contains 63 electrons and has an electron configuration of [Xe]4f76s2. For Eu2+, the electron configuration is [Xe]4f7. Due to the divalent oxidation state, the lowest excited levels originate from the 4f 65d1 configuration, and thus, energy absorption leads to the promotion of one 4f electron to the 5d orbital. However, in the trivalent oxidation state, only six electrons remain in the 4f shell, leaving one 4f orbital unoccupied in terms of Hund’s rules. As the nuclear charges of trivalent Ln ions are less screened than those of divalent Ln ions, the energies of the 5d orbitals of trivalent Ln ions are typically higher in the vacuum ultraviolet (VUV) range [22–26]. Thus, these ions dominantly show narrowline 4fn ↔ 4fn transitions in the UV-Vis range [27]. Emission spectra related to transitions between the 5d and 4f orbitals of Ln ions, such as for Eu2+ and Ce3+, are characterised by broad bands, whose properties are highly dependent on the coordination environment of the ions. The 4fn−15d1 → 4fnrelated emission bands can vary between the near-UV and the near-IR [28] spectral ranges depending on the covalence and crystal field conditions around the electrons in the Ln 5d orbitals. For instance, the emission peak maximum shifts to higher wavelengths when Eu2+ is doped into a highly covalent host lattice with soft anions, including sulfides [29] or nitrides [30], due to the so-called nephelauxetic effect. The nephelauxetic effect is a ligand field effect that stabilises the energies of 5d orbitals by the delocalisation of the electrons over the contributing orbitals of the ligands. Metal–ligand (ML) distance also influences the emission energy. A short ML distance causes a large splitting of the 5d orbital and consequently a shift to longer wavelengths. Electric dipole (ED)-allowed 4fn−15d1 → 4fn transitions, observed for divalent and trivalent Ln ions, result in up to six orders of magnitude higher emission intensities than those in the cases of electric-dipole-forbidden transitions. Owing to the half-filled 4f7 shell, Eu2+ is substantially stable, as indicated by its low standard potential in aqueous solution (E0(Eu3+/Eu2+) = −0.35 V) and is thus suitable for technological and industrial applications. In contrast, other divalent Ln ions are not very stable and are difficult to handle under ambient conditions. For Ln ions such as Sm3+, Eu3+, Tb3+, and Dy3+, whose photoluminescence typically occurs in the visible spectral range, radiative relaxation from excited 4fn levels are observed. These transitions are spin- and parity-forbidden and are specific for each ion. For example, orange, red, and green lights are characteristics of the emissions of the 4f orbitals of Sm3+, Eu3+, and Tb3+, respectively. As the 4f electrons are shielded by the outer 5s and 5p electrons, the positions of the spectral peaks associated with these

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electronic orbitals are not influenced by the environment of the ions and remain nearly unchanged for various host lattices [27, 31]. Owing to the low variation in the emission properties of trivalent Ln ions because of the shielding of the 4f electrons, Dieke et al. [32] and Carnall et al. systematically investigated the energy levels involved in the luminescence of La ions in LaCl3 [22] and LaF3, respectively [33]. By doping LaCl3 with different trivalent Ln ions and using a parameterised crystal field model, semi-empirical determination of the energy levels involved for each unique ion was conducted, and the acquired values are summarised in the Dieke diagram. Emission spectra of trivalent Ln ions are characterised by sharp peaks covering the energy ranges typical for each ion. The main emission peaks of Eu3+ and Tb3+ are located in the red (5D0 →7 FJ, J ≤ 3) and green spectral ranges (5D4 →7FJ, J ≤ 6), respectively. With the Dieke diagram, the electronic transitions for an observed emission peak can be assigned. Extension of the energy levels to the VUV range was reported by Wegh et al. [23] Among the ions that undergo the 4fn → 4fn electronic transitions, Eu3+ has the highest potential for use as a coordination sensor for monitoring structural changes and the mechanisms of chemical reactions. Effects of the symmetry of the coordination environment on the photoluminescence spectra are comprehensively discussed in this section. Eu3+ is often applied ex situ as a spectroscopic probe to obtain information about the surrounding coordination sphere [31, 34]. A primary reason for this is the lack of degeneracy of both the ground 7F0 and lowest emitting 5D0 levels, which considerably facilitates the interpretation of the luminescence spectra. Elucidation of the spectra is also promoted by the low total angular momentum number (J) values of the most dominant emissive 5D0 → 7FJ transitions (J = 0–2) and the differences between the energies of the different 7FJ levels that allow clean separation, and thus, assignment of the respective luminescence peaks [31]. Moreover, the crystal field potential exhibited by the locally surrounding ligands of Ln ions further removes the degeneracy of the 2S+1LJ levels depending on the local symmetry. The lack of degeneracy of the 5D0 and 7F0 levels and the correspondingly low number of arising spectral lines facilitate the application of Eu3+ as an ex situ probe for crystal field symmetry analysis [31]. The total number of microstates in the electronic configuration [Xe]4f6 of Eu3+ is calculated to be 3,003 using the following equation: ! 14 14! . (5:1) = n!ð14 − nÞ! n Interelectronic repulsion, spin–orbit coupling, crystal field splitting, and Zeeman effect are the key parameters determining the energetic position and are responsible for lifting the degeneracy of Eu3+ microstates. Spin–orbit coupling is a relativistic effect and may be best considered as a result of the interactions between the spin magnetic moments of electrons and the orbital magnetic moments arising from the movement of

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electrons around the nucleus, thereby mixing the orbital and spin angular momenta, respectively. Crystal field splitting is a special case of the Stark effect that additionally removes the MJ degeneracy of the levels owing to the presence of an electric field. However, the Zeeman effect leads to the same consequence in the presence of a magnetic field. Both external fields explicitly break the symmetry of the observed system, thereby further splitting the different spin–orbit levels. Nevertheless, in the case of an odd number of electrons, time-reversal symmetry must be considered, which requires the presence of Kramers doublets even in a crystal field. In contrast, the Zeeman effect destroys the time-reversal symmetry [27]. Generally, the electronic levels of Ln ions originating upon the inclusion of the mutual electron repulsion and spin–orbit coupling are determined in the limit of the Russell–Saunders coupling scheme that enables the designation of the electronic levels by the 2S+1LJ term symbols. In 2S+1LJ, S is the total spin quantum number leading to 2S + 1 degenerate MS states. The number 2S + 1 is referred to as spin multiplicity and is denoted as singlet (1), doublet (2), triplet (3), and so on. L is the total orbital angular momentum quantum number and is designated by the letters S, P, D, F, G, H, I, K, L, M . . . for L = 0, 1, 2, 3, 4, 5, 6, 7, 8, 9 . . ., respectively. J is presented by the different couplings between L and S in the vector model, that is: J = L + S, L + S − 1, L + S − 2, . . ., jL − Sj

(5:2)

The previously described Russell–Saunders coupling scheme is applicable in the case of stronger mutual electron repulsion than spin–orbit coupling. Nevertheless, for Ln ions, spin–orbit coupling is already non-negligible because of their high atomic numbers, and an intermediate coupling scheme would be more appropriate. Consequently, the designation of the electronic levels by term symbols loses its meaning because strong spin–orbit coupling mixes L with S. However, the Russell–Saunders scheme still offers a reasonable starting point for a qualitative description of the energy levels of Ln, as presented in the Dieke diagram; nevertheless, this description should only be regarded as a crude approximation [27]. As a representative example, we consider the ground levels of Eu3+. Eu3+ has a 4f 6 ground configuration. The term with the highest spin multiplicity is a septet as S = 6 × 1/2 = 3, and consequently, 2S + 1 = 7. According to Hund’s rules, L may be determined by the sum of all ml values of the occupied orbitals, which affords: L = ½ð+3Þ + ð+2Þ + ð+1Þ + 0 + ð−1Þ + ð−2Þ = 3

(5:3)

and thus leads to an F term based on the designation rules. The separation between different 2S+1L terms is approximately 10,000 cm−1 for the lower terms of the 4f 6 configuration of Eu3+ (Figure 5.3, right panel) [31]. These terms further split into 295 2S+1LJ levels due to spin–orbit interaction in the case of Eu3+ [27]. For the 7F Eu3+ term, when S = 3 and L = 3, the possible values for J are 6, 5, 4, 3, 2, 1, and 0, resulting in the well-known 7FJ (J = 0–6) ground levels of Eu3+. The respective degeneracy in MJ of each spin–orbit level is 2 J + 1. The abovementioned shielding effects

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of the 5s and 5p electrons from the coordinative environment on the 4f electrons lead to approximately fixed energy positions of the 7FJ and higher excited 5DJ’ (J’ = 0, 1, 2) levels, irrespective of the host compound. However, the crystal field splitting of the 2J + 1 degenerate spin–orbit levels is of the order of ≤102 cm−1 (Figure 5.3) [31]. Crystal field perturbation destroys the spherical symmetries of free ions, and the 2S+1 LJ levels split into several crystal field levels depending on the site symmetry of Ln3+. In the case of Eu3+, the lack of degeneracy of both the lowest excited 5D0 and 7F0 ground levels simplifies the situation such that the corresponding splitting pattern of the other J levels may directly provide information about the local symmetry of the Ln ion in a surrounding ligand field (Table 5.1 and Figure 5.4). Table 7.1 shows the splitting of the different J levels into the corresponding Stark states based on the local point symmetry of Eu3+. For example, in Oh symmetry, only one peak was noticed for the 5D0 →7F1 transition, and no peaks were observed for the other transitions. This is primarily because of transition selection rules. According to the Laporte selection rule, intraconfigurational ED transitions (including s–s, p–p, d–d, and f–f electronic transitions) are forbidden as they result in vanishing transition matrix elements. However, this rule is only strictly applicable to Ln ions in the gas phase [31]. To understand the different numbers of possible emission peaks for the 5D0 → 7FJ (J = 0, . . ., 6) transition in a certain point group, it is necessary to initially elucidate the mechanisms of Laporte-forbidden 4fn ↔ 4fn transitions in detail followed by the effects of symmetry reduction on these transitions. A majority of the Eu3+ transitions are induced by ED transitions. This means that the Eu3+ transitions occur due to the interaction between the time-averaged ED moment in an Ln ion caused by the 4f electrons and the electric field vector of the external light field. As the ED moment is proportional to the position vector, its sign changes under an inversion operation, that is, it has an odd parity. Thus, any transition matrix elements between 4f orbitals should also have overall odd parities in a centrosymmetric environment and will accordingly vanish. Nevertheless, these selection rules can be avoided by admixing the 4f orbitals with other orbitals of different parities or coupling to vibrations. The first mechanism is the basis of the Judd–Ofelt mechanism, which is briefly described below [31]. In addition to ED transitions, orbital and spin angular momenta induce magnetic moments that may couple to the magnetic field component of light, thereby leading to magnetic dipole (MD) transitions. Unlike ED transitions, MD transitions are allowed by the Laporte selection rule because of the even parity of the magnetic moment; however, their intensities are weaker by a factor of approximately α2 ~ 10–4 as compared to those of the ED transitions, where α is the fine structure constant. Therefore, MD transitions become relevant if the ED transitions are suppressed. In the case of Eu3+, the 5D0 → 7F1 transition is an important example of an MD transition, which is dominant in centrosymmetric site symmetries. Note that the 4f electrons in Ln may even induce quadrupole moments that can couple to the electric field gradient of light. The line strengths of ED transitions are of the same order of magnitude as those of the MD transitions [31].

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Figure 5.3: Emission spectrum of [Eu(tta)3(phen)] (tta = 2-thenoyltrifluoroacetylacetonate, phen = 1,10-phenanthroline, T = 77 K, λex = 396 nm) showing transitions from the 5D0 state (left) (reproduced with permission from reference [31]). Energy diagram of Eu3+ (4f6) with relative magnitude of the interelectronic repulsion (terms), spin–orbit coupling (levels), and the crystal field potential (sublevels) (reprinted with permission from reference [58]. Copyright 1987 Elsevier).

A breakthrough in the understanding of the occurrence of ED-type 4fn ↔ 4fn transitions was independently achieved by Judd and Ofelt in the 1960s, respectively [35, 36]. Their key assumption was the admixture of energetically higher lying 5d orbitals in the 4fn configuration and the fact that the energy difference was substantially larger than that between different J levels of the original 4fn configuration. Under these conditions, an effective ED moment operator with even parity (including the odd part of the crystal field potential) could be derived with perturbation theory. This framework readily explains the induced ED transitions in Ln. However, note that the Judd–Ofelt theory only enables intensity modelling of the absorption and luminescence spectra by a leastsquares fitting procedure to the experimental spectrum and does not predict the entire spectrum. Nevertheless, to date, it is the only theoretical model that completely describes the emergence of ED-type 4fn ↔ 4fn transitions. Moreover, it explains the generally observed trend in the Eu3+ spectra that the 5D0 → 7FJ transitions with J = 2, 4, and 6 are typically stronger than those with odd J values, owing to new selection rules that can be derived from group theory in the theoretical frameworks of these transitions. An exception is centrosymmetric crystal fields that only contain even components in the multipole expansion, for which the Judd–Ofelt theory is not appropriate because of the failure of one of its assumptions [31]. The 5D0 → 7F0 transition of Eu3+ is strictly forbidden in terms of both common and Judd–Ofelt selection rules. However, its occurrence provides extremely valuable information because both levels with J = 0 do not split in the crystal field. Thus, the splitting of this transition directly indicates the occupancy of more than one crystallographically independent site in a specific host, revealing, for instance, the presence of side products

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or reaction intermediates. If the structural difference between these sites is small, the presence of two or more different species can be indicated by an asymmetric shape, a shoulder, or broadening of the emission peak assigned to the 5D0→7F0 transitions. Luminescence spectra of single sites can be acquired by site-selective excitation experiments. There are two main reasons for the emergence of the peak of this transition in the luminescence spectra. On the one hand, odd crystal fields facilitate the mixing of other terms with higher J values (which is termed as J mixing) [37, 38]. Therefore, the appearance of the 5D0 ↔ 7F0 emission peak at approximately 577–581 nm implies that Eu3+ typically occupies a low-symmetric site, such as Cnv, Cn, or Cs symmetry site. Another mechanistic explanation for the emergence of this peak was suggested by Wybourne and Downer et al. [39–41] They considered the effects of configuration interaction on a third-order ED moment mediated by 4f–5d spin–orbit mixing, which significantly contributes to this transition. However, for Eu3+, J mixing is probably the most dominant contribution due to the large energy difference between the 4f6 and 4f55d1 configuration levels [42, 43]. As abovementioned, the intensity of the emission peak at 585–600 nm, attributed to the 5D0 → 7F1 transition, is substantially independent of the coordination environment of Eu3+. Nevertheless, it can be also employed to acquire structural information. For instance, for comparison of different emission spectra, the intensities can be scaled such that the 5D0 → 7F1 peaks have the same intensity in both spectra. In addition, the number of emission peaks observed for this emissive transition provides information about the crystal field symmetry: cubic and icosahedral crystal fields lead to no splitting; hexagonal, trigonal, and tetragonal crystal fields cause a splitting into two lines; and orthorhombic and lower symmetry crystal fields result in a splitting into three lines. As abovementioned, this is related to the successive loss of the J degeneracy [31]. Emission peaks of the 5D0 → 7F2 transition are located between 610 and 625 nm. This hypersensitive transition is strictly forbidden if an inversion centre is present at the site occupied by Eu3+. According to the Judd–Ofelt mechanism, this transition should have a very low intensity when Oh symmetry sites are occupied. SrTiO3 with a cubic perovskite structure is an illustrative example of this case. When SrTiO3 is doped with Eu3+, a negligibly small emission peak of the 5D0 → 7F2 transition is achieved due to the centrosymmetric 12-fold-coordinated Sr2+ site with Oh symmetry. Owing to this property, the ratio between the intensities of the peaks of 5D0 → 7F2 and 5D0 → 7F1 transitions can be used to measure the order of symmetry of the Eu3+ site [31]. The 5D0 → 7F3 transition is a Judd–Ofelt-forbidden transition, and the intensities of the emission peaks at 640–655 nm corresponding to this transition are very weak. Intensities of the emission peaks of 5D0 → 7F4 electronic transition at 680–710 nm must be carefully evaluated, because the spectra must be appropriately corrected owing to the low sensitivities of photomultiplier tubes (PMTs) in this spectral range. As the emission peaks of the 5D0 → 7F5 (740–770 nm) and 5D0 → 7F6 (810–840 nm) transitions are located beyond the detection limits of most UV-Vis PMT detectors, these transitions are not comprehensively discussed in many articles [31].

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In addition to the analyses of the splitting behaviours and intensity ratios of the peaks of Eu3+ transitions, further information about structural changes in the coordination environments of cation sites can be obtained by investigating additional features, including “solvent quenching” and “antenna effect”, of the 4fn → 4fn transitions [27]. Solvent quenching refers to the strong quenching of the emission intensities of Ln ions, such as Eu3+, induced by energy transfer to OH valence vibrations when these ions are submerged in protic solvents including water and ethanol, both of which are regularly used in solution-based syntheses of solid compounds [44]. As Ln ions exhibit intense quenching effects in solution, significant enhancements in their emission intensities are expected during the precipitation of Ln-containing compounds as the solvent molecules in the coordination sphere around the Ln ions are gradually exchanged by ligands.

Figure 5.4: Flowchart indicating the number of emission peaks observed for the 5D0 → 7FJ (J = 1, 2, 4, 6) transitions of Eu3+ depending on the site symmetry of Eu3+. σ denotes polarisation of the emitted light parallel to the incident electric field, and π represents polarisation of the emitted light perpendicular to the incident electric field (reprinted with permission from reference [59]).

If these ligands are organic, precipitation often results in the formation of complexes, coordination polymers, or metal–organic frameworks (MOFs), which often boost the emission intensities of Ln3+ via the antenna effect. Although trivalent Ln ions have very attractive emission properties, their absorption cross sections are significantly low. The combination of Ln centres with organic ligands facilitates the transfer of the excitation energy absorbed by, for example extended conjugated π-systems of the organic ligands to Ln ions, thereby enhancing the electronic excitations and improving the luminescence intensities of these ions [45].

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Table 5.1: Number of emission peaks observed for the 5D0 → 7F0-4 transitions of Eu3+ for different point group symmetries [30]. Point group C Cs C Cv Ci Ch D Dh Dd D C Cv Ch Ci Dd Dh C Ch Cv Dh Dd S D C Cv D Ch Dh T Td Th O Oh Ih



D, F                                  



F                                  



F                                  



F                                  



F                                  

5.4 Examples of multimodal in situ experimental instrumentation As mentioned earlier in Section 5.2, an important feature of ILACS is the application of an external reactor in it, which enables data acquisition under real reaction conditions. Due to its flexible stationary and mobile setups, this approach can be combined with different commercial synthesis workstations [2, 17]. These reactor systems enable

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the control of synthesis parameters, such as stirring, dosing, and temperature, during the monitoring of reaction parameters, including pH, ion conductivity, redox potential, and concentrations of ions (e.g. Ca2+, Cu2+, Na+, K+, Cl–, and NO3–). Figure 5.5 shows the examples of experimental setups for measuring in situ luminescence using an excitation source in reflection (Figure 5.5, top) and transmission (Figure 5.5, bottom left) modes. A Y-shaped optical fibre is employed to simultaneously transmit excitation light from the spectrometer to the reactor solution and the emitted light from the reactor to the detector. The transmission excitation mode is generally used for a mobile in situ setup comprising a transportable luminescence spectrometer, equipped with a CCD detector and an optical fibre. Excitation energy is provided by external UV light-emitting diodes (LEDs), and luminescence is evaluated in the transmission mode using an optical fibre submerged in the reactor. This example of

Figure 5.5: Examples of the experimental setups for multimodal in situ measurements. Top: A Fluorolog 3 spectrometer equipped with a Y-shaped optical fibre for excitation in the reflection mode. Bottom left: UV LED excitation, allowing simultaneous monitoring of in situ pH and ion conductivity. Bottom right: Integrated in situ luminescence, IR spectroscopy, and XRD setup at the DESY PETRA III facility (adapted from references [2, 48] with permission from the Royal Society of Chemistry).

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mobile setups can be transported to other laboratories such as synchrotron facilities, which is advantageous, because it facilitates the use of simultaneous complementary analyses, for example, in situ luminescence and XRD. In situ luminescence measurements in the transmission mode are also useful because the detection of the extinction of excitation light can afford additional information about the turbidity of a solution. However, the reactant concentration must be kept low to prevent complete blockage of the excitation light. In contrast, in situ luminescence measurements in the reflection mode are less sensitive to the reactant concentration. In detail, to conduct simultaneous in situ luminescence and XRD experiments at the synchrotron facilities, the reactor has to be slightly modified to decrease the pathway by which the synchrotron radiation must transverse through the solvent volume. One example of a modified reactor vessel possessing an indented glass tube through the reaction vessel is shown in Figure 5.6. The reactor holder accordingly developed for this vessel has been designed to quickly fit different beamlines and does not require realignment after each reaction. It contains two openings for the entrance and exit of X-ray beams in addition to a larger opening arranged at 90° to the X-ray path for UV-LED irradiation during luminescence measurements and is combined with integrated stirring and heating systems [46]. Generally, mobile spectrometer systems are employed in synchrotron applications [2]. Alternatively, a few beamlines offer integrated luminescence measurements within their standard infrastructure.

Figure 5.6: Adapted glass reactor for simultaneous in situ luminescence and XRD measurements (left) and the reactor holder (right) (reprinted with permission from reference [46]).

5.5 Selected application examples Due to their high emission intensities, high crystallinity, and simplified syntheses, luminescent complexes are ideal model systems for demonstrating the proof-of-principle

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of monitoring the formation of functional materials using in situ luminescence measurements. Therefore, the first practical examples commented in this Section are based on different 1,10-phenanthroline (phen) and 2,2ʹ-bipyridine (bipy)-containing luminescent Ln complexes [2, 47, 48]. Luminescent complexes are in general relevant, for example, technological and biomedical applications such as for the production of organic LEDs (OLEDs) and cell markers for bioimaging. For instance, in situ luminescence measurements can be applied for monitoring metal–ligand exchange processes during the complexation in ethanol and decomposition of the complex [Eu(phen)2(NO3)3] with the addition of Sn4+. In a simplified exemplary experiment [2], a solution of the phen ligand was introduced into the reactor containing a solution of europium(III) nitrate. As expected, before the addition of the phen molecules, the emission spectrum of Eu3+ in ethanol had very low intensity and very low resolution for the 5D0 → 7F0-4 electronic transitions of Eu3+, owing to the strong quenching effect of the OH vibrations of solvent molecules. Upon addition of phen, the emission intensity significantly increased and its spectral resolution improved, indicating the gradual exchange of the solvent molecules in the coordination sphere of Eu3+ for the phen ligand and the initiation of an antenna effect. Nevertheless, this increase in luminescence intensity occurred 3–5 min after the introduction of the phen molecules, implying delayed nucleation of [Eu(phen)2(NO3)3]. These results were confirmed by simultaneous synchrotron-based in situ XRD and in situ IR spectroscopy as well as via complementary in situ measurements of pH and ion conductivity [2]. Formation of [Eu(phen)2(NO3)3] was also detected by the strong shift of the peaks corresponding to the 5D0 → 7F4 transition of Eu3+ to higher energies. This blue shift reveals a larger energy gap between the 5D0 excited state and the 7F4 ground state, probably because of the stronger interaction between Eu3+ and the phen ligands than that between Eu3+ and the coordinating ethanol molecules of the solvation shell. This stronger interaction is expected due to the slight nephelauxetic effect observed in N-containing solid materials with high stability and short interatomic distances. Upon the dissolution of SnCl2 in the ethanol solution and the oxidation of Sn2+ to Sn4+, the phen ligands detached from Eu3+ and coordinated Sn4+, producing [Sn(phen)Cl4]. [Sn(phen)Cl4] is non-luminescent in the whole regarded UV-Vis spectral range at the applied excitation wavelength of 395 nm, and its formation is detected by a considerable decrease in the emission intensity [2]. Polzin et al. [47] used in situ luminescence measurements to investigate the influences of temperature and reactant concentration on the induction time and crystallisation rate, on the formation of reaction intermediates and their stability, as well as on the ligand exchange process induced by the introduction of phen ligands to the [Eu(bipy)2 (NO3)3] complex. The results revealed that heating the reaction system from room temperature to 50 °C reduced the induction time, indicated by an earlier increase in the spectral signals of the 5D0 → 7F1 transition. These results were confirmed by the earlier growth of the Bragg reflections of [Eu(bipy)2(NO3)3] during in situ XRD. A further increase in temperature to 65 °C did not significantly influence the reaction kinetics [47].

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It was additionally demonstrated that when the reactant concentration was reduced by 50%, the onset of crystallisation increased from 7 to 13 min (Figure 5.7, left). When the reactant concentration was further decreased to 40% of that used in the initial experiments, the crystallisation rate decreased; furthermore, when the reactant concentration was decreased to 30%, the crystallisation onset was delayed to t = 28 min, and the crystallisation rate further slowed down. These in situ luminescence results enabled the calculation of the function describing and predicting the relationship between the crystallisation time and the reactant concentration. Analysis of the acquired function showed that the crystallisation time exponentially increased with a decrease in the reactant concentration and enhanced towards infinite values when the reactant concentration was below 30%. These predictions were verified by decreasing the reactant concentration to 10%, resulting in no product generation. Interestingly, although the ion concentrations were low, spectral emissions of Eu3+ could still be detected, allowing Polzin et al. to estimate the ILACS detection limit, which was 2 mM Eu3+ in solution [47]. Moreover, during the crystallisation of [Eu(bipy)2(NO3)3], an increase in the reactant concentration to 300% did not only reduce the induction time to t ≈ 1.5 min, but also altered the reaction mechanism inducing the formation of a reaction intermediate prior to product formation. The presence of intermediates could be deduced from the changes in the rate of emission intensity in the spectral data. Compared to the smooth exponential increase observed in the intensity of the 5D0 → 7F2 transition at 593 nm at lower reactant concentrations (Figure 5.7, left), the intensity of the 5D0 → 7F2 transition at 593 nm at higher reactant concentrations initially increased and then exhibited a prolonged plateau period (approximately 9.5 min) before increasing again until t ≈ 17 min and at the end of the reaction (λex = 365 nm; Figure 5.7, right) [47]. Analysis of the splitting pattern of the 5D0 → 7FJ emission peaks confirmed the formation of the reaction intermediate as it demonstrated that at t = 0, 5 and 20 min, Eu3+ in the reaction solution was exposed to three distinct coordination environments. At t = 0 min, Eu3+ was only coordinated by the solvent molecules, and the final product was verified to be fully formed at t = 20 min. The definite splitting pattern at t = 5 min strongly suggested the generation of a reaction intermediate. Examination of the emission spectrum at t = 5 min provided important insights into the nature of this intermediate structure. For example, the single peak of the 5D0 → 7F0 transition indicated that only one symmetry-equivalent site could be occupied by Eu3+ in this compound. High intensity of the peak of the 5D0 → 7F2 transition implied that Eu3+ did not occupy centrosymmetric sites, whereas the two resolved peaks assigned to the 5D0 → 7F1 transition indicated an intermediately high crystalline symmetry probably with three-, four-, or six-fold local site symmetries in a tetragonal or hexagonal crystal system. Interestingly, although the concomitant in situ XRD measurements confirmed the formation of an intermediate via the increase of additional reflections at t = 5 min, the high background and broad reflections in the XRD data prevented the elucidation of any structural information. Even ex situ XRD failed to generate useful structural data for this compound as the intermediate was highly unstable.

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Figure 5.7: In situ emission spectra (λex = 365 nm) showing changes in the crystallisation rates of [Eu(bipy)2(NO3)3] in response to changes in the reactant concentration (left). Detection of the formation of a reaction intermediate during the crystallisation of [Eu(bipy)2(NO3)3], confirmed by concomitant synchrotron-based in situ XRD experiments (reproduced from reference [47] with permission from the PCCP Owner Societies).

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A final feature of this study was a set of experiments in which phen was added after the production of [Eu(bipy)2(NO3)3] to monitor the incorporation of phen molecules into the Eu3+ compound. This process can be easily detected by in situ luminescence via a significant increase in emission intensity. The increase in emission intensity was expected due to the presence of an additional aromatic ring in the phen ligand structure, which offered an enhanced antenna effect. Similar to the previous examples, this phase transition was also verified by in situ XRD [47]. Although the antenna effect on the 4f 6 → 4f 6 electronic transitions in Eu3+ is very useful for examining the incorporation of organic ligands into the coordination sphere of Eu3+, it still provides limited information about the uptake of Eu3+ in solution for the formation of a solid material during actual precipitation. This is because the positions of the peaks assigned to the 4fn → 4fn transitions are not significantly influenced by changes in the crystal field of Eu3+. In contrast, the 4fn−15d1 → 4fn electronic transitions in Eu2+ (n = 7) and Ce3+ (n = 1) exhibit a substantially stronger influence of the crystal field splitting owing to the interaction of the excited state 5d electron with its local surrounding. The emission spectra of these electronic transitions are strongly shifted by complexation, the covalencies of binding ligands (nephelauxetic effect), lengths of newly formed bonds, and changes in the coordination number. As Eu2+ easily oxidises to the more stable Eu3+ (which lacks 4f 55d → 4f 6 transitions in the visible range) under ambient conditions, Ce3+ is a more appropriate ion for use as a coordination sensor [27]. Accordingly, Arana et al. [48] used Ce3+ as a coordination sensor to monitor the crystallisation of the model system [Ce(phen)2(NO3)3]. The emission spectrum of Ce3+ in ethanol (λex = 400 nm) showed a broad 5d1 → 4f1 Ce3+-based transition band before the addition of the ligand, with a maximum at approximately 545 nm and Commission Internationale d’Éclairage (CIE) 1931 colour coordinates of x = 0.4494, y = 0.4339 (Figure 5.8). After product formation, the emission band (λex = 400 nm) shifted, with a maximum at approximately 660 nm and the CIE 1931 colour coordinates of x = 0.6693, y = 0.3231. This considerable difference between the optical properties of Ce3+ in the solvation shell and those of Ce3+ in [Ce(phen)2(NO3)3] facilitated the tracking of the reaction progress by in situ luminescence measurements. To enable simultaneous monitoring of both the reactant and the product while minimising the band overlap, the low-energy emission band of Ce3+ in the product was monitored at 700 nm, and the intensity of the high-energy emission of Ce3+ as a reactant was examined at 500 nm. Figure 5.8 shows a comparison of the emission intensities at 500 and 700 nm, the volume of the added phen solution, and the in situ measurement results of pH and ion conductivity. To appropriately analyse the mechanism of the formation of [Ce(phen)2(NO3)3] quantitatively, the process was divided into three stages: A, B, and C (Figure 5.8). Reaction times of the stage A, B, and C were 0–3, 3–10, and >10 min, respectively [48]. Stage A consisted of the induction time. In that stage, the pH increased due to the addition of the alkaline phen solution; nevertheless, the intensity of the luminescence bands at 500 and 700 nm and the ion conductivity did not change significantly, indicating

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that the product was not yet formed. In contrast, in stage B, crystallisation began, as revealed by the increase in the emission intensity of the luminescence band at 700 nm and the simultaneous decrease in the emission intensity of the luminescence band at 500 nm. These results were confirmed by the concomitant changes in the pH and ion conductivity of the reaction solution. At approximately t = 3 min, the pH started to decrease owing to the uptake of the phen molecules from the solution followed by their incorporation into the solid complex, confirming product formation. Thereafter, the pH increased again until t = 10 min because of the additionally added ligand molecules. In this stage, ionic conductivity also began to decrease at t = 3 min owing to the uptake of free ions from the solution for the product formation, which further confirmed the onset of crystallisation. In stage C, the addition of the ligand solution was discontinued, and therefore, the pH stopped increasing. Instead, it decreased until approximately t ≈ 11 min and then stabilised. The same behaviour was observed for the ion conductivity, which decreased up to t ≈ 11 min and subsequently stabilised. Similarly, no further increase in the emission intensity of the band at 700 nm or decrease in the emission intensity of the band at 500 nm was detected during stage C, implying no considerable complex

Figure 5.8: Emission spectra (top left) and CIE 1931 chromaticity diagram (bottom left) of Ce3+ before and after the formation [Ce(phen)2(NO3)3] in ethanol. Comparison of the in situ luminescence results with the time dependence of the addition of phen to Ce3+-containing solutions, pH and ion conductivity during the crystallisation of [Ce(phen)2(NO3)3] (reproduced from reference [48] with permission from the Royal Society of Chemistry).

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formation after this period and the completion of the reaction. The increase in the emission intensity of the band at 500 nm during the stages B and C was probably caused by a slight overlap of this band with the low-energy emission band [48]. As explained in Section 5.1, in situ luminescence measurements can be employed to examine the production of non-emissive materials when coordination sensors are introduced into their structure. To demonstrate this principle, Eu3+ and Ce3+ substituted Ca2+ ions in selected calcium phosphates to monitor the formation of and transformations between different calcium phosphate phases. Calcium phosphates are medically highly relevant because it is the main inorganic constituent of mammals and is extensively used in the production of implants and prostheses. Understanding the transformation between different calcium phosphate phases can help deduce the conversions between healthy and pathological tissues, prevent and treat diseases, and improve the stability of implants. In the abovementioned example, in situ luminescence was used to detect the production of amorphous calcium phosphate (ACP) and Ca5(PO4)3OH (hydroxyapatite, HAp), besides their phase transformations to CaHPO4 · 2H2O and Ca8H2(PO4)6 · 5H2O under real reaction conditions [49]. Firstly, to investigate the influence of the concentration of coordination sensors on the reaction pathway, preliminary experiments were performed to compare the phase transformations of undoped calcium phosphates and the ones doped with 3, 5, and 7 mol% Ce3+. Results of these preliminary experiments showed that, undoped, a mixture of ACP and HAp was initially formed. After 7 min, these phases converted to CaHPO4 · 2H2O, and upon heating to 80 °C, an additional transition to Ca8H2(PO4)6 · 5H2O occurred (t = 50 min). Doping of 3 mol% Ce3+ into the Ca2+ sites did not affect phase transitions in this system. However, when 5 mol% Ce3+ was doped into the Ca2+ sites, a transformation of ACP to HAp occurred, followed by the conversion of HAp to CaHPO4 · 2H2O after t ≈ 10 min. When 7 mol% Ce3+ was doped into the Ca2+ sites, CaHPO4 formed prior to the crystallisation of CaHPO4 · 2H2O. In addition, introduction of 7 mol% Ce3+ into the Ca2+ sites enhanced amorphisation, increasing the length of the stable ACP phase, possibly due to the increase in disorder caused by the differences between the charge and ionic radii of the two cations. To analyse the influence of the type of coordination sensor (Ce3+ or Eu3+) on the phase transformations of calcium phosphate, the investigated compounds were doped with both ions, and their powder diffractograms were compared. Introduction of Ce3+ or Eu3+ into the host lattice led to the same phase transformations at the same reaction times. In conclusion, 3 mol % of Ln ions or the type of coordination sensor (namely, Ce3+ and Eu3+) did not affect the phase transformations of calcium phosphates. Therefore, 3 mol% of both coordination sensors were employed to further examine the respective phase transformations via in situ luminescence analysis [49]. For this purpose, a mixed aqueous solution of Ca(NO3)2 and Ce(NO3)3 was added to a (NH4)2HPO4 solution. As expected, no luminescence was detected for the (NH4)2HPO4 solution; however, within 10 min of the addition of the cation solution, a broad emission band between 310 and 440 nm with a maximum at 353 nm emerged, indicating product

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formation (Figure 5.9). This emission range is consistent with the electronic transitions of Ce3+ from the lowest 5d1-based states to the 2F5/2 and 2F7/2 ground state levels of Ce3+ in HAp [50]. The asymmetrical shape of the Ce3+-based emission band implied the presence of a secondary phase, which was confirmed by ex situ XRD. At t = 13 min, the emission intensity started to considerably decrease, suggesting the production of a new phase with lower emission intensity. Ex situ XRD demonstrated the formation of CaHPO4 · 2H2O during this stage of the reaction, which explained the decrease in emission intensity, because this structure introduced quenching H2O molecules into the coordination sphere of Ce3+. Although the in situ emission spectra were recorded at a constant excitation wavelength of 280 nm, the decrease in the emission intensity could have alternatively been caused by a shift of the excitation bands during the phase transformations of calcium phosphates. To verify this hypothesis, the excitation spectra were additionally recorded for the abovementioned reaction. Timedependent intensity of the excitation spectra presented the same profile as the timedependent intensity of the emission spectra, and no substantial shift was observed in the excitation bands. This indicated that the variation in the intensity of the excitation spectra nearly exclusively correlated with the variation of the corresponding emission properties [49]. CaHPO4 · 2H2O formation was also revealed by a slight redshift of the emission spectra probably caused by the higher coordination number and shorter average bond lengths of CaHPO4 · 2H2O as compared to those of HAp. As ACP lacks long-range order, limited structural information is available about this phase [51–53]. At t = 20–40 min, the luminescence intensity remained constant, indicating no significant structural changes in this time range. After t = 40 min, the temperature started to increase and the emission intensity started to decrease due to thermal quenching. At t ≈ 51 min, the emission intensity slightly increased and then decreased again, implying an additional structural change in the cation environment and a phase transformation at T ≈ 60 °C. In summary, ex situ and synchrotron-based in situ XRD confirmed the conversion of CaHPO4 · 2H2O to Ca8H2(PO4)6 · 5H2O during this reaction stage. The production of Ca8H2(PO4)6 · 5H2O was also revealed by the in situ recorded emission spectrum via the emergence of an additional emission band, a slight redshift and intensity increase, expected owing to the increase in the number of distinct crystallographic sites for the doping of Ce3+ and a decrease in the number of coordinated water molecules in Ca8H2(PO4)6 · 5H2O. A similar behaviour was also detected using Eu3+ as a coordination sensor for tracking the same reaction pathway [49].

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Figure 5.9: Transformations between different calcium phosphate phases studied by in situ luminescence measurements (left). In situ emission spectra recorded during the crystallisation of calcium phosphate using Ce3+ as a coordination sensor (reproduced from reference [49] with permission from the Chinese Chemical Society (CCS), Peking University (PKU), and the Royal Society of Chemistry).

5.6 Tracking the reaction pathways without coordination sensors In general, ILACS setups can be additionally used to track formation of solid materials that do not involve coordination sensors. This goal is achieved by monitoring changes of other emitting species such as organic molecules or monitoring the solution turbidity by measuring the time-dependent intensity of the light transmitted through the solvent during chemical reactions [46]. Measurements of light transmission are especially advantageous because the time for formation of solid materials can be estimated irrespective of the luminescence properties and is nearly independent of the wavelength of the light source; this increases the flexibility of its application. In the subsequent example, a 395-nm LED was used as the light source. Thus, the light source irradiated the reactor from the outside, whereas the CCD detector connected to an optical fibre submerged in the reactor continuously measured the intensity of the light transmitted through the reaction medium. A combination of in situ luminescence analysis and light transmission was applied, for instance, to examine the formation and phase transformations of [Al(acac)3] (acac = acetylacetone) [46], an important precursor for the synthesis of Al2O3 nanomaterials [54]. In summary, for the synthesis of [Al(acac)3], an aqueous solution of NH3 was added to an aqueous solution containing Al(NO3)3 · 9H2O and acac. In situ measurements helped determine the synthesis conditions resulting in proper crystallisation of [Al(acac)3] and the influence of the concentration of NH3 in the reaction solution (that is, pH) on the induction time. In daylight, although [Al(acac)3] is colourless, it emits blue light under UV irradiation (320 nm, Figure 5.10). Ex situ luminescence studies

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indicate that the emission of [Al(acac)3] (λex = 320 nm) ranges from 400 to 700 nm and consists of multiple bands peaking at 440, 460, and 509 nm. In contrast, pure acac shows very weak luminescence with an emission band between approximately 417 and 833 nm (maximum, 550 nm) upon excitation at λex = 380 nm [46]. Crystallisation can be examined, therefore, via in situ luminescence measurements because of the large differences between the optical properties of uncoordinated acac molecules and those of the acac molecules as ligands coordinating Al3+ in [Al(acac)3] [46]. Thus, the progress of the synthesis of [Al(acac)3] was characterised via in situ luminescence analysis by monitoring the intensity of the blue emission band at 450 nm under excitation at 320 nm. Initially concentrated NH3 (25%) was added to an aqueous solution of Al(NO3)3 · 9H2O and acac. After approximately 2.1 min, product formation was detected by a substantial increase in the emission intensity upon the addition of 1.06 mL or ~ 14 mmol of NH3, and the critical pH of ~ 2.4 was reached. At t = 3.2 min, the emission intensity started to decrease, continuously decreasing until the end of the reaction, with the exception of an oscillation at approximately t = 20–30 min. In situ light transmission measurements and synchrotron-based in situ XRD confirmed the formation of [Al(acac)3], its partial dissolution after t ≈ 3.2 min, and its conversion from α- and γ-phases at approximately t = 20–30 min, justifying the changes detected by the in situ luminescence measurements in this time range. When the NH3 concentration was decreased to 12.5%, an onset of the emission intensity occurred only after 3.7 min, which suggested a delay in product formation to the reaction time when the amount of NH3 approached the critical concentration and the pH reached ~ 2.5. When the NH3 concentration was decreased to approximately 3.5%, the emission intensity increased only after 11.2 min and the pH reached 2.5, verifying the successful control of product formation via the adjustment of the NH3 concentration [46].

Figure 5.10: Blue emission of [Al(acac)3] under UV light (left). Time-dependent emission intensity of [Al (acac)3] at 450 nm (λex = 320 nm) showing the influence of NH3 concentration on the crystallisation time of [Al(acac)3] (right) (reproduced with permission from reference [46]).

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5.7 Conclusions and outlook This chapter provides a brief introduction to the new ILACS technique, which was recently developed at the Kiel University. It explores the theoretical background and advantages, experimental setup, and current applications. Although in situ luminescence analysis shows high potential for monitoring and elucidating the mechanisms of chemical reactions in many settings, this approach is still in its infancy and needs to be further developed. To date, this strategy has mostly focused on examining the formation of solid materials via simplified co-precipitation synthesis techniques; nevertheless, in future, it can be also potentially used to investigate solvothermal, ultrasound, microwave, ionic liquid, solid-state, or melting-based syntheses. For example, to facilitate the application of in situ luminescence analysis for these synthesis methods, practical and technological adaptations to physical equipment are necessary. Solid-state or melting-based syntheses can, for instance, be easily monitored exploring a reflection geometry of emission and excitation radiation and by using an observation window within the applied oven. Many new experimental studies can be reached adapting the in situ luminescence approach to flow reactor cells, which is currently in development. In general, expansion of integrated experimental setups that enable concomitant monitoring of reactions by combining additional approaches such as pair distribution function (PDF) analysis is also one additional goal to widen the scope and practicality of ILACS. In addition to new synthesis and analysis techniques monitoring the formation of additional types of materials, including quantum dots and metal-organic frameworks, can be combined with in situ luminescence measurements. Furthermore, manipulating the reaction conditions to allow the regular use of Eu2+ as a coordination sensor owing to its intense and coordination-sensitive 4f 65d1 → 4f7 emission bands may open many new areas of study for new types of materials. Finally, the spectroscopic changes expected during the conversion between the different oxidation states of Ce and Eu ions can be also used for in situ monitoring of redox processes.

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[23] R. T. Wegh, A. Meijerink, R.-J. Lamminmäki, J. Hölsä, Extending Dieke's diagram. J. Lumin. 87–89, 1002–1004, 2000. [24] L. Van Pieterson, R. T. Wegh, A. Meijerink, M. F. Reid, Emission spectra and trends for 4fn−15d ↔ 4fn transitions of lanthanide ions: Experiment and theory. J. Chem. Phys. 115, 9382–9392, 2001. [25] L. van Pieterson, M. F. Reid, R. T. Wegh, S. Soverna, A. Meijerink, 4fn→4fn−15d transitions of the light lanthanides: Experiment and theory. Phys. Rev. B 65, 045113/1–16, 2002. [26] L. van Pieterson, M. F. Reid, G. Burdick, A. Meijerink, 4fn→4fn−15d transitions of the heavy lanthanides: Experiment and theory. Phys. Rev. B 65, 045114/1–13, 2002. [27] W. M. Yen, H. Yamamoto, S. Shionoya, Phosphor Handbook. Boca Raton: CRC Press Laser and Optical Science and Technology, 2006. [28] N. Kunkel, H. Kohlmann, A. Sayed, M. Springborg, Alkaline-earth metal hydrides as novel host lattices for Eu-II luminescence. Inorg. Chem. 50, 5873–5875, 2011. [29] D. Jia, X. Wang, Alkali earth sulfide phosphors doped with Eu2+ and Ce3+ for LEDs. Opt. Mater. 30, 375–379, 2007. [30] P. Pust, P. J. Schmidt, W. Schnick, A revolution in lighting. Nat. Mater. 14, 454–458, 2015. [31] K. Binnemans, Interpretation of europium(III) spectra. Coord. Chem. Rev. 295, 1–45, 2015. [32] G. H. Dieke, Spectra and Energy Levels of RE Ions in Crystals. New York: Wiley, 1968. [33] W. T. Carnall, G. L. Goodman, K. Rajnak, R. S. Rana, A systematic analysis of the spectra of the lanthanides doped into single-crystal LaF3. J. Chem. Phys. 90, 3443–3457, 1989. [34] K. Binnemans, C. Görller-Walrand, Application of the Eu3+ ion for site symmetry determination. J. Rare Earths 14, 173–180, 1996. [35] B. Judd, Optical absorption intensities of rare-earth ions. Phys. Rev. 127, 750–761, 1962. [36] G. Ofelt, Intensities of crystal spectra of rare‐earth ions. J. Chem. Phys. 37, 511–520, 1962. [37] G. Nishimura, T. Kushida, Local field in glass probed by laser-induced fluorescence-line narrowing in Ca(PO3)2:Eu3+. Phys. Rev. B 37, 9075–9078, 1988. [38] M. Tanaka, G. Nishimura, T. Kushida, Contribution of J mixing to the 5D0−7F0 transition of Eu3+ ions in several host matrices. Phys. Rev. B 49, 16917–16925, 1994. [39] B. G. Wybourne, Effective operators and spectroscopic properties. J. Chem. Phys. 48, 2596–2611, 1968. [40] M. C. Downer, G. W. Burdick, D. K. Sardar, A new contribution to spin‐forbidden rare earth optical transition intensities: Gd3+ and Eu3+. J. Chem. Phys. 89, 1787–1797, 1988. [41] G. W. Burdick, M. C. Downer, D. K. Sardar, A new contribution to spin‐forbidden rare earth optical transition intensities: Analysis of all trivalent lanthanides. J. Chem. Phys. 91, 1511–1520, 1989. [42] M. Tanaka, T. Kushida, Interference between Judd-Ofelt and Wybourne-Downer mechanisms in the 5 D0−7FJ (J=2,4) transitions of Sm2+ in solids. Phys. Rev. B 53, 588–593, 1996. [43] M. Tanaka, T. Kushida, Optical transition mechanisms of Sm2+ and Eu3+ in solids. J. Lumin. 72–74, 141–143, 1997. [44] A. Beeby, I. M. Clarkson, R. S. Dickins, S. Faulkner, D. Parker, L. Royle, A. S. De Sousa, J. A. Gareth Williams, M. Woods, Non-radiative deactivation of the excited states of europium, terbium and ytterbium complexes by proximate energy-matched OH, NH and CH oscillators: an improved luminescence method for establishing solution hydration states. J. Chem. Soc., Perkin Trans. 2, 493–504, 1999. [45] K. Binnemans, Lanthanide-Based Luminescent Hybrid Materials. Chem. Rev. 109, 4283–4374, 2009. [46] N. Pienack, L. Ruiz Arana, W. Bensch, H. Terraschke, In situ studies on phase transitions of Tris(acetylacetonato)-aluminum(III) Al(acac)3. Crystals 6, 157/1–14, 2016. [47] P. Polzin, I. V. Eliani, J. Ströh, M. Braun, N. Ruser, N. Heidenreich, P. Rönfeldt, F. Bertram, C. Näther, S. Wöhlbrandt, M. Suta, H. Terraschke, From ligand exchange to reaction intermediates: What does really happen during the synthesis of emissive complexes? Phys. Chem. Chem. Phys. 20, 7428–7437, 2018.

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[48] L. Ruiz Arana, P. Lindenberg, H. Said, M. Radke, N. Heidenreich, C. D. Santos Cunha, S. Leubner, H. Terraschke, Monitoring the mechanism of formation of [Ce(1,10-phenanthroline)2(NO3)3] by in situ luminescence analysis of 5d-4f electronic transitions. RSC. Adv. 7, 52794–52800, 2017. [49] H. Terraschke, M. Rothe, A.-M. Tsirigoni, P. Lindenberg, L. Ruiz Arana, N. Heidenreich, F. Bertram, M. Etter, In-situ luminescence analysis: A new light on monitoring calcium phosphate phase transitions. Inorg. Chem. Front. 4, 1157–1165, 2017. [50] I. V. Berezovskaya, N. P. Efryushina, E. V. Zubar, V. P. Dotsenko, “Distribution and Luminescent Properties of Ce3+ Ions in Nanosized Calcium Hydroxyapatite,” Sumy State University, p. 01PCN24,1 2012. [51] S. V. Dorozhkin, Calcium orthophosphates in nature, biology and medicine. Materials 2, 399–498, 2009, and references therein. [52] A. S. Posner, F. Betts, N. C. Blumenthal, Formation and structure of synthetic and bone hydroxyapatites. Progr. Cryst. Growth Char. 3, 49–64, 1980. [53] J. E. Harries, D. W. L. Hukins, C. Holt, S. S. Hasnain, Conversion of amorphous calcium phosphate into hydroxyapatite investigated by EXAFS spectroscopy. J. Cryst. Growth 84, 563–570, 1987. [54] L. D. Hafshejani, S. Tangsir, H. Koponen, J. Riikonen, T. Karhunen, U. Tapper, V.-P. Lehto, H. Moazed, A. A. Naseri, A. Hooshmand, Synthesis and characterization of Al2O3 nanoparticles by flame spray pyrolysis (FSP) – Role of Fe ions in the precursor. Powder Technol. 298, 42–49, 2016.

Huayna Terraschke✶

Chapter 6 In situ monitoring of the syntheses, phase transformations, and loading processes of metal–organic frameworks Metal–organic frameworks (MOFs) are formed via the “LEGO”-like self-assembly of ligand and metal building blocks. However, in situ measurements during MOF formation have shown that the prediction of MOF crystal structures is not simple. In practice, minor variations in the synthesis parameters can disturb the gradual organisation of the building blocks into the desired ordered crystalline compounds, leading to the formation of different polymorphs or even non-porous structures. This chapter summarises important discoveries in this field and newly developed reactors for probing the chain of events leading to MOF formation and the loading of guest molecules within MOF pores.

6.1 Introduction As explained in Chapter 2, metal–organic frameworks (MOFs) are compounds formed by cationic inorganic building units (IBUs) interconnected by polydentate anionic or neutral organic ligands. IBUs are typically isolated ions or metal oxo compounds, whereas common organic ligands include amines, carboxylates, phosphonates, and azolates [1]. This combination of inorganic and organic components yields highly crystalline materials with extended networks, unique porosity, and high specific surface areas. Unlike other porous materials, such as periodic mesoporous organosilicas and zeolites, MOFs exhibit highly tuneable architectures and chemical properties [2]. Even minor changes in the structure, length, and functional groups of the organic ligands or the composition of the metal salts can drastically alter the nature of the resulting MOF. Nevertheless, various examples have been reported in which the chemical and structural properties were deliberately varied, with the aim of obtaining MOFs that are potentially useful for a wide range of applications, including the separation of gaseous fuels (e.g. H2), CO2 capture and storage, catalysis, drug delivery, light harvesting, light production, and molecular sensing [3]. Although MOF formation occurs spontaneously via self-assembly, in most cases, the complexity of the chemical reaction prevents detailed predictions about the structure of the final MOF. The gradual reorganisation of disordered solution species into an ✶

Corresponding author: Huayna Terraschke, E-Mail: [email protected]

https://doi.org/10.1515/9783110459098-006

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ordered crystalline structure involves many individual physical and chemical interactions, each of which depends on multiple synthesis parameters, such as temperature, reactant concentration, and reaction time [4]. Due to the large number of reaction parameters and the cumulative nature of MOF formation, even small changes in the synthesis conditions can result in the formation of very different MOF polymorphs or even completely different compounds [5]. The lack of mechanistic understanding is a serious obstacle [3] for the development of new pre-designed MOFs and the transfer of synthesis procedures from the laboratory to large-scale production, which is necessary for MOF commercialisation [6]. For these reasons, obtaining detailed information about the influence of synthesis parameters on the crystallisation mechanism of MOFs is essential to control their structure formation and structure-related properties, especially the porosity. To elucidate the pathway of MOF formation, analyses must be performed not only after synthesis, or ex situ, but also during the synthesis process using so-called in situ analysis techniques. In situ techniques can provide extremely detailed information about crystallisation processes, phase transformations, synthesis kinetics, side products, and MOF stability under real synthesis conditions. This chapter is focused on two main topics: the in situ monitoring of MOF synthesis and phase transformations and the in situ analysis of MOF loading processes. In particular, Section 6.2 reviews selected examples from the current literature on the characterisation of MOF formation and phase transformations during solution and solvent-free synthesis processes, especially using in situ techniques such as synchrotron-based X-ray diffraction (XRD) analysis and in situ infrared (IR) spectroscopy. In situ pair distribution function (PDF) measurements, which have only recently been applied as a promising method for monitoring MOF formation, are also summarised in this chapter. In addition, examples for the use of in situ characterisation methods for solution-based reactions in both batch and flow reactors are also reviewed. Sections 6.3 and 6.4 focus on the analysis of MOF loading processes with guest molecules in the bulk and in thin films, respectively, in which the quartz crystal microbalance (QCM) technique plays a key role. The growing recognition of the usefulness of in situ techniques for analysing MOF formation combined with the improved functionality provided by third-generation synchrotron sources and resolution-enhanced detectors [7] has led to rapid growth in the number of MOF analyses employing in situ monitoring, over the last several years. Hence, the total volume of publications in this field cannot be reviewed in a single chapter. Therefore, the following sections focus on selected examples of scientific articles (mostly published between 2015 and 2020) that specifically address the analysis of MOF formation, phase transformations, and loading processes. For further information about studies on the in situ analysis of MOF synthesis published prior to 2015, please see the reviews by Walton and Millange [8] and Férey et al. [9]. For mechanistic insights into the construction of reactors for in situ studies, the reader should refer to the work of O’Hare et al. [10], Walton and O’Hare [11], and Terraschke et al. [12] For an overview on the quantitative kinetic modelling of MOF

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formation, the review by Van Vleet et al. [13] is recommended. The reader may also be interested in the article by Reinsch and Stock [14], which provides a personal view on the rationalisation, application, and exploration of MOF synthesis, and the paper by Heinke et al. [15], which summarises recent developments in the synthesis and in situ characterisation of surface-mounted MOFs (SURMOFs).

6.2 In situ analysis of MOF synthesis and phase transformations 6.2.1 In situ XRD analysis of solution-based processes In situ XRD analysis employing batch reactors As the most valuable tool for monitoring changes during the formation or transformation of crystalline MOFs, in situ XRD provides detailed and high-resolution data. However, this technique is not without challenges. For successful in situ XRD measurements, the high background signal produced by the reactor walls and bulk solvent in solutionbased synthesis processes is a significant technological hurdle that must be overcome to obtain acceptable signal-to-noise ratios and analysable data sets. This issue can be addressed by using high-energy, high-flux X-rays (white or monochromatic); such X-rays are available at synchrotron facilities, such as the Deutsches Elektronen-Synchrotron (DESY) in Germany, the European Synchrotron Radiation Facility (ESRF) in France, the Source Optimisée de Lumière d’Energie Intermédiaire du LURE (SOLEIL) in France, the Diamond Light Source in the UK, and MAX IV in Sweden, to name only a few. While these facilities provide access to the necessary high-intensity X-rays, they are removed from the normal laboratory setting and lack the equipment required to perform on-site MOF synthesis. Standard beamline infrastructure generally does not include the equipment required for MOF synthesis because of the extreme reaction parameters used, such as high pressures, high temperatures, ultrasonic or microwave energy, and specific pH levels, viscosity levels, reactant concentrations, and solvent compositions [12]. Several research groups have overcome these technological challenges, and this section summarises successful experiments, the technological solutions employed, and the obtained insights into MOF chemistry. Since 2015, multiple in situ XRD studies on MOFs have been facilitated by the development of a new, versatile, and easy-to-use remote-controlled reactor set-up by the Kiel University and the DESY facility (Figure 6.1, left). This set-up, named SynRAC (synchrotron-based reaction cell for the analysis of chemical reactions), has several key features that facilitate in situ XRD measurements. In addition to integrated openings for the Xray beam, the reactor cell includes a stirring unit, a heating mantle, and injection systems for liquids and solids. The heating mantle can control the temperature within the

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Figure 6.1: Schematic representation of the SynRAC set-up developed at the Kiel University in cooperation with DESY (left). Reaction progress (α) and nucleation probability (Pn) as a function of time for the synthesis of Ce-UiO-66 at 70, 80, 90, and 100 °C (right) (reprinted from reference [7], with the permission of AIP Publishing).

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reaction vessel in the range from ambient room temperature to 180 °C and is cooled with compressed air. The injection systems are capable of accurately adding either two liquid reactants or a single solid reactant to the reaction at pre-set intervals. To simplify use, each SynRAC unit is designed to hold two sizes of commercially available disposable borosilicate vessels (5 and 11 mL) [7]. Examples of studies that have employed SynRAC cells to examine MOF formation include Heidenreich et al. [7, 16], Reinsch et al. [17], Rhauderwiek et al. [18], and Bueken et al. [5], each of which is discussed below. To verify the performance of the different features of the SynRAC reactor, various test reactions were carried out at beamline P08 [19] at DESY [7]. First, the crystallisation kinetics of Ce-UiO-66 ([Ce6(OH)4(O)4(BDC)6], where UiO = University of Oslo and BDC = benzene-1,4-dicarboxylate) were investigated by analysing the obtained in situ XRD data using the Gualtieri approach [20]. The progress of reactions performed at 70, 80, 90, and 100 °C was fitted with a non-linear equation (eq. 6.1) based on the Gualtieri model [20] (Figure 6.1, right) [7]: α=

  
1   · 1 − exp kg · t n t−a 1 + exp − b

(6:1)

where α is the reaction progress, t is the reaction time, kg is the rate constant of crystal growth, n is the dimension of growth, b is the distribution of the probability of nucleation with time, and a is inversely proportional to the nucleation rate constant, kn [7]: a=

1 . kn

(6:2)

The probability of nucleation, Pn, can also be calculated using the following equation (Figure 6.1, right) [7]: ! −ðt − aÞ2 Pn = exp (6:3) b In addition, plotting ln(k) as a function of 1/T gave the activation energies for nucleation (50(5) kJ/mol) and crystal growth (32(5) kJ/mol) [7]. Following these proof-of-concept studies, the SynRAC reactor was employed to synthesise the coordination polymer (CP) [Bi(HIDC)(IDC)] (H2IDC = 4,5-imidazoledicarboxylic acid). This CP evolved from a precursor phase composed of the highly disordered bismuth nitrate compound [Bi6O4(OH)4]0.54[Bi6O5(OH)3]0.46(NO3)5.54, which was formed as a crystalline intermediate [7, 21]. Moreover, employing citraconic acid as a linker and a green-chemistry-based approach with mild reaction conditions and water as the only solvent, the in situ XRD study showed the formation of a novel MOF, CAU-15-Cit ([Al2 (OH)4(O2C-C3H4-CO2)] · nH2O, where CAU = Christian-Albrechts-Universität) [16]. In addition, SynRAC was used in studies examining the green synthesis of the MOF Al-MIL-68-Mes ([Al(OH)(O2C-C3H4-CO2)] · nH2O), where MIL represents Matériaux

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de l’Institut Lavoisier), under mild reaction conditions [17]. AI-MIL-68-Mes consists of cationic aluminium centres interlinked by mesaconate (MES) and hydroxide ions. Al-MOFs are of commercial interest because aluminium salts are inexpensive, abundant, and they generate crystalline compounds with high structural stability, even at extreme temperatures [17]. The efficient synthesis of Al-MIL-68-Mes, which would facilitate its widespread industrial use, requires a detailed understanding of the crystallisation process. For this reason, the formation of AI-MIL-68-Mes was analysed by in situ XRD at beamline P09 [22] at DESY. These studies revealed the immediate formation of an amorphous phase upon mixing of the reactants, followed by the formation of a highly crystalline phase. As the Gualtieri approach was not appropriate for fitting the experimental curves, the reaction kinetics were evaluated via a Sharp–Hancock plot based on the Avrami theory [23]. Specifically, ln(−ln(1 − α)) was plotted as a function of ln(t − t0) for the time period t0 to t, during which a constant reaction mechanism dominated the crystallisation process. Although this type of analysis is restricted to evaluating the experimental data obtained from homogenous systems such as solutions or gels, it can be effectively employed to provide general information about the reaction dynamics of nearly uniform systems. The slope of the linear regression model allowed for the determination of the Avrami exponent for the reaction. This metric, which is related to the rate-limiting step for product formation, indicated that nucleation was the rate-limiting step for the crystallisation process. Thus, the crystallisation of Al-MIL-68-Mes consumed the reactants remaining in solution that were not consumed by the precipitation of the intermediate amorphous phase. The plot intercept, n·ln(k), could also be used to estimate a pseudo rate constant for the reaction [17]. Rhauderwiek et al. [18] reported four novel Ce3+-porphyrin-based MOFs (PMOFs) named CAU-18, CAU-18a, CAU-19-X, and Ce-PMOF-4NO2. The crystallisation processes of CAU-18 ([Ce4(H2TCPP)3(DMF)2(H2O)4]) and CAU-19-H ([Ce3(H2TCPP)2(BA-X)(HBA-X/H2O)2] · 2HBAX · nH2O, where H6TCPP is 4-tetracarboxyphenylporphyrin, DMF is dimethylformamide, X = H, 2Cl, 3Cl, 4Cl, 3CO2H, 4NH2, or 4NO2, and HBA = C7H4O2) were studied using in situ XRD at beamlines P07B [24] and P09 [22] at DESY. Employing the analysis method developed by Gualtieri [20], the obtained in situ XRD data were used to determine the nucleation probability (Pn) and the Arrhenius activation energies of both nucleation (kn) and crystal growth (kg). The sensitivity of the kinetic parameters, specifically, the induction time, conversion time, and nucleation probability, to the synthesis temperature was also investigated, and an inverse relationship was revealed between the reaction temperature and each kinetic parameter [18]. Detailed information about the stability and transformation kinetics of UTSA-74 ([Zn2(DOBDC)], (where DOBDC is 2,5-dioxidobenzene-1,4-dicarboxylate and UTSA = University of Texas at San Antonio) into its polymorph Zn-MOF-74 in water was obtained using the SynRAC reactor [5]. For these experiments, UTSA-74 crystals were suspended in deionised water and heated to either 110 or 120 °C in the SynRAC reactor, and timeresolved in situ XRD (Δt = 30 s) data were collected over 40 min at beamlines P08 [19] (λ = 0.51662 Å) and P09 [22] (λ = 0.53905 Å) at DESY. In situ XRD analysis (Figure 6.2) revealed the simultaneous appearance of reflections assigned to Zn-MOF-74 and the

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progressive loss of the UTSA-74 signal intensity, beginning at minute 6 for the reaction at 110 °C and at minute 5 for the reaction at 120 °C. In contrast to previous reports, these results demonstrated that USTA-74 is unstable in the presence of water at both temperatures [25] and transforms via a dissolution–recrystallisation process.

Figure 6.2: In situ XRD data (λ = 0.51662 Å) showing the conversion between UTSA-74 and Zn-MOF-74 at (a) 110 °C and (b) 120 °C (republished with permission of Royal Society of Chemistry, from reference [5]; permission conveyed through Copyright Clearance Center, Inc.).

The crystallisation of two three-dimensional MOF polymorphs, SION-1 and SION-2 ([Tb2 (DHBDC)3DMF)4] · 2DMF, where DHBDC = 2,5-dihydroxybenzene-1,4-dicarboxylate, SION refers to MOFs synthesized at the EPFL Valais in Sion, Switzerland) developed at the École Polytechnique Fédérale de Lausanne (EPFL) Valais in Sion (Switzerland), was studied using in situ XRD at the Swiss-Norwegian beamline (SNBL) at ESRF [4]. The reaction of Tb(NO3)3 · 6H2O with H2DHDBC in a DMF:H2O solvent mixture at 120 °C for 24 h yielded the golden SION-2 polymorph. Further reaction for an additional 48 h resulted in the conversion of SION-2 into the bright red SION-1 polymorph. More detailed information about the processes involved in the conversion of SION-2 into SION-1 was obtained from a second set of XRD experiments in which pre-formed SION-2 immersed in DMF:H2O was slowly heated (2 °C/min) to 120 °C and held at this high temperature for 6 h. Time-resolved measurements (Δt = 60 s) revealed that SION-2 was partially dissolved at high temperatures, which allowed conversion to the more thermodynamically stable SION-1 polymorph. Moorhouse et al. [26] employed in situ diffraction experiments to analyse the reaction mechanism during the resin-assisted solvothermal synthesis of [Co(NDC)(DMF)] (NDC is 2,6-naphthalenedicarboxylate). For this work, the Oxford-Diamond in situ cell (ODISC) furnace, initially developed at beamline I12 at the Diamond Light Source [27], was used at HASYLAB beamline F3 at DORIS (DESY) for energy-dispersive X-ray diffraction (EDXRD) measurements (white beam energy = 15 ≤ E/keV ≤ 65, detector angle (2θ) =

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1.4°). This analysis required the use of high-energy in situ XRD to penetrate the ODISC. In the resin-assisted solvothermal synthesis, cation-impregnated macroporous resin beads (typically commercially available sulfonated styrene-divinylbenzene pre-treated to allow adsorption of the metal cation of interest) fused to an amorphous, cross-linked, infinite network of interpenetrating polymer chains are used as both a template and an ion source for the reaction. This approach provides enhanced control over cation release in solution. Heating the reaction solution containing metal-impregnated resin beads to mild or moderate temperatures ( 0.292 (Figure 6.14). When the x-value was decreased (x = 0.5–2), the lower concentration of the BME-BDC ligand and higher concentration of the DB-BDC ligand caused the transition pressure to shift to a higher value. In contrast, for x = 0, the compound underwent a guest-independent heat-induced np–lp transition upon activation, which also influenced the adsorption behaviour [72]. Detailed studies combining single-crystal XRD and in situ XRD at the Pohang Accelerator Laboratory in Korea were able to identify the nature of gate-opening processes in a flexible (flex) MOF upon CO2 adsorption. In this study, the flex MOF contained a macrocyclic Ni2+ complex ([Ni(C14H34N6)]2+, ([NiLpropyl]2+) and a 2,2,5,5-biphenyltetracarboxylate (BPTC) linker (Figure 6.15), resulting in a composition of {[(NiLpropyl)2(BPTC)] · 4DMF · 4H2O}. When activated, the flex MOF exhibited a closed structure, with the three pendant arms of its macrocycles blocking the pore apertures. However, in the presence of CO2, the macrocycles rotated along the axis linking the carboxylate molecules to the Ni2+ coordination centre, allowing CO2 molecules to access the pores (Figure 6.15) [70]. Interestingly, this flex MOF was also found to exhibit breathing behaviour. CO2 adsorption in the MOF pores expanded the rectangular dimensions of their openings from 9.19 × 20.78 Å2 to 12.00 × 29.79 Å2 (Figure 6.15). This expansion was primarily attributed to an increase in the dihedral angle between the phenyl rings, carboxylate planes, and benzene rings of the BPTC ligands upon CO2 adsorption [70]. An additional MOF example whose gate-opening mechanism has been well studied is Ni-DUT-8 ([Ni2(NDC)2(DABCO)], DUT = Dresden University of Technology). Singlecrystal XRD, in situ XRD, and extended X-ray absorption fine structure (EXAFS) analyses revealed that this compound exhibited an unprecedented 254% increase in pore

Chapter 6 In situ monitoring of the MOF syntheses and phase transformations

Figure 6.14: In situ adsorption of CO2 (195 K) (left) and corresponding in situ XRD patterns (right) for [Zn2(BME-BDC)x(DB-BDC)2−xDABCO]n (x = 2) (reprinted from publication reference [72]. Copyright 2015, with permission from Elsevier).

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size upon the adsorption of gases, which was far larger than the volume change previously observed for any other gate-opening MOF. Interestingly, the extent of pore deformation depended on the adsorption enthalpy as well as the adsorbate shape and kinetic diameter. In fact, Ni-DUT-8 was very sensitive to adsorbates because different gaseous probe molecules generated very different absorption kinetics. For example, the adsorption of nitrogen (77 K), carbon dioxide (195 K), and n-butane (272.5 K) induced one-step structural transformations from closed to large pores, whereas the adsorption of ethane (185 K) and ethylene (169 K) induced a two-step transformation [71]. Other MOFs, such as Al-MIL-53 doped with V4+ ions, exhibit structural breathing behaviour in response to temperature, pressure, and air humidity. This breathing phenomenon was monitored by in situ electron paramagnetic resonance spectroscopy by Nevjestić et al. [73].

Figure 6.15: (a) Macrocyclic complex ([Ni(C14H34N6)]2+, ([NiLpropyl]2+) and H4BPTC linker components of the flex MOF, (b) Gate-opening, and (c) breathing behaviour of the flex MOF upon CO2 adsorption (reprinted with permission from reference [70]. Copyright 2016, American Chemical Society).

6.3.4 Negative gas adsorption (NGA) transitions In conventional isothermal gas adsorption experiments, the gas uptake is expected to increase with pressure, but the anomalous negative gas adsorption [74] effect can cause the unexpected spontaneous desorption of gas upon increasing the pressure. One important example of an extensive study combining in situ XRD, gas adsorption, and EXAFS experiments at the BESSY II light source (Helmholtz-Zentrum Berlin) during the adsorption of methane and n-butane by DUT-49 ([Cu2(BBCDC)], where BBCDC = 9,9′-([1,1′-biphenyl]-4,4′-diyl)bis(9H-carbazole-3,6-dicarboxylate) showed that this phenomenon was triggered by structural phase transformations of the framework. DUT-49 [75] exhibited an open-pore structure (DUT-49op) consisting of an assembly of cuboctahedral

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metal–organic polyhedra, which were composed of copper paddlewheel units and BBCDC ligands. This ordered arrangement resulted in the formation of large open pores (10, 17, and 24 Å). During NGA, DUT-49op underwent contraction, with the unit cell volume decreasing considerably from 100,071.8(5) Å3 to 47,282.0(3) Å3 for the closed-pore structure (DUT-49cp). This transition was reversible and DUT-49cp transformed back to DUT-49op via an intermediate pore structure (DUT-49ip) once the pores were completely filled. These structural changes were also reflected by the space group symmetry and lattice parameters of the op, cp, and ip phases of DUT-49 [74].

6.4 Loading processes on MOF thin films MOF thin films are potentially important for applications in membranes for separation or sensing. SURMOFs, an important example of thin-film MOFs, are often produced by a layer-by-layer technique employing liquid-phase epitaxy. In general, these materials are synthesised by the alternating submersion of a functionalised substrate in solutions of metal ions and organic linkers, with the number of cycles determining the film thickness. The application of MOFs as sensors often depends on their functionalisation by loading with guest molecules (e.g. luminescent molecules). To obtain detailed information about the mass transfer and diffusion properties during the incorporation of guest molecules, the loading processes can be monitored in situ using the QCM technique [15]. A QCM generates acoustic waves and measures the changes in the wave propagation properties during adsorption and desorption processes. QCMs are piezoelectricbased devices that can detect small changes in electrode mass, with a sensitivity of up to 10 μg/m2 [76]. The Sauerbrey equation establishes a linear relationship between the resonant frequency and small mass increments, as follows [76]: 2 ΔM Δf0 = − pffiffiffiffiffiffiffiffi · f02 · A E·ρ

(6:4)

where f0 is the fundamental resonant frequency obtained when only odd overtones are excited, Δf0 is the variation in the resonant frequency, E is the stiffness (Young’s modulus), ρ is the density of the propagating material, ΔM is the mass increase, and A is the area of the sensor surface. This equation is applied under the assumption that the mass deposited at the surface follows the vibration of the crystal, and the sensor becomes thicker [76]. This method has been applied to study interfacial processes on surfaces and thin films, e.g. species transport, adsorption kinetics, and film growth. A QCM-based study for analysing loading processes was performed during the adsorption of europium β-diketonate compounds within the pores of HKUST-1, which can act as a photonic antenna for Eu3+ excitation [60]. Luminescence spectroscopy after the loading procedure

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revealed successful doping. An increase in mass of only 160 ng/cm2 rather than the expected mass increase of 500 ng/cm2 was detected, which indicated that only approximately one in three super-cage pores housed a Eu3+ complex after the loading procedure (Figure 6.16) [60]. The QCM results gave a diffusion constant of ~1 × 10–19 m2/s, which is reasonable, because the [Eu(BZAC)3BIPY] (BZAC = benzoylacetonate, BIPY = 2,2-bipyridine) molecules have a large cross-section (0.97 nm) relative to the size of the entrance of the channels connecting the MOF pores (1.2 nm). The low diffusion rate could also be caused by the pore entrances being blocked by [Eu(BZAC)3BIPY] molecules crystallised on the MOF surface. However, this hypothesis was ruled out because only reflections assigned to the MOF thin film were detected in the XRD measurements (Figure 6.16) [60]. 180 (400) 160 (200)

140 Intensity / a.u.

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40 20 0 0

500

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1500 2000 Time / s

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8

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Figure 6.16: QCM results for the incorporation of [Eu(BZAC)3BIPY] guest molecules on a 40-layer HKUST-1 SURMOF (left). XRD patterns of the HKUST-1 SURMOF before (black curve) and after (grey curve) the uptake of [Eu(BZAC)3BIPY]. Inset: laterally patterned HKUST-1 SURMOF sample under UV light, in which the bright features correspond to SURMOF structures loaded with [Eu(BZAC)3BIPY] (republished with permission of Wiley, from reference [60]; permission conveyed through Copyright Clearance Center, Inc.).

The QCM technique was also used to monitor the formation of Bi2O3 semiconductor nanoparticles (NPs) within the pores of a HKUST-1 SURMOF. The high porosity, regular pore size, and high loading capacity of this MOF provided a favourable platform for producing monodispersed NPs. The resulting Bi2O3 NPs showed enhanced photoefficiency as compared to conventionally synthesised NPs for the photodegradation of a dye (nuclear fast red). The diameters of the Bi2O3 particles (1–3 nm), as revealed by transmission electron microscopy, were slightly larger than the inner diameter of the HKUST-1 pores (1.9 nm), which was explained by the authors by the displacement of the BTC ligand molecules defining the MOF pore walls during particle formation [77].

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6.5 Concluding remarks In summary, this chapter provided an overview of selected examples (published mostly between 2015 and 2020) on the in situ monitoring of MOF formation and loading processes, with the aim of enabling a better understanding of these phenomena. Due to the large number of reported in situ diffraction experiments, this chapter focused on studies that applied synchrotron X-ray radiation, which is necessary to penetrate the reactor walls and liquid media commonly used for MOF synthesis. In recent years, formation mechanisms have been most commonly studied using ZIF, HKUST-1, UiO, and CAU MOFs. Interestingly, the crystallisation processes of ZIFs have been monitored using in situ UV/Vis absorption spectroscopy [46], PDF measurements [50], light scattering [53, 54], and XRD analysis in flow reactors [35] or during mechanosynthesis [6]. The formation of HKUST-1 has been also studied by in situ diffraction experiments during mechanosynthesis [38], IR spectroscopy [42, 45], and AFM [56]. UiO MOFs have been intensively investigated during synthesis using real-time PDF measurements [48, 49] and XRD in batch [7] or flow [37] reactors. For a great variety of CAU MOFs (CAU-8 [34], CAU15 [16], CAU-18 [18], and CAU-19 [18]), XRD analysis has not only been performed during solvothermal synthesis but also during the adsorption of guest molecules, such as hydrogen within CAU-1 [59]. Fascinating in situ diffraction experiments for monitoring the adsorption of other gases, including carbon dioxide in [66] and oxygen [67], have revealed remarkable structural changes such as the breathing effect in CAU-13 [1] and MIL-53 [73] or NGA within DUT-49 [74]. Finally, this chapter also demonstrated the applicability of time-dependent QCM measurements for monitoring the loading of guest molecules within MOF thin films, such as luminescent materials within a HKUST-1 SURMOF [60].

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Huayna Terraschke✶, Katia Nchimi, Felix Hartmann, Wolfgang Bensch

Chapter 7 Operando studies on the charge and discharge processes of battery materials To supply sustainable, renewable energy, the efficiency of its storage, such as within batteries, is as decisive as the efficiency of its production. To produce powerful rechargeable batteries, the development of electrodes with low cost, long cycle and calendar life, high current density, and improved safety is crucial. Operando studies during the charge and discharge processes enable a profound understanding of how electrode materials work and why they fail. This chapter offers a short overview of the application of frequently used operando techniques for characterising the function of lithium-, sodium-, potassium-, zinc-, and magnesium-ion batteries with liquid or solid electrolytes.

7.1 Introduction In the recent decades, significant progress has been made in the generation of renewable energy using, for example, sunlight or wind as a power source. Compared to fossil fuels such as gas, oil, or charcoal, these sources of energy offer distinct advantages such as their potentially unlimited nature and the absence of produced greenhouse gases. However, owing to the periodicity of their occurrence, these renewable energies are not permanently available, despite their relatively high abundance. Moreover, unlike fossil fuels, these sources usually cannot serve directly as primary energy sources because they cannot be transported. Sustainable use requires their storage in secondary devices that will supply energy even in the absence of a primary energy source, leading to stable and reliable clean power grids [1]. In addition to the improvement of existing electrical networks, the introduction of batteries for energy storage is particularly important for the automotive industry. Electric vehicles offer the advantage of potentially maintaining a high level of mobility without releasing dangerous gases such as carbon and nitrogen oxides into the atmosphere [2]. Thus, the utilisation of renewable energy depends not only on the ability to collect and transform energy but also on the ability to store the transformed energy, for example, in batteries. Thus, the availability and reliability of batteries are serious considerations in popularising sustainable energy forms.

✶

Corresponding author: Huayna Terraschke, E-Mail: [email protected]

https://doi.org/10.1515/9783110459098-007

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To harness renewable energy, efforts are being made towards the development of battery materials offering low cost, long-term cycling, high current density, and safety for large-scale energy storage and mobile (e.g. vehicular) applications. In this context, the performance of rechargeable batteries depends not only on their chemical composition but also on the stability of their components, especially the electrodes. Optimally, the charge and discharge processes should be accompanied by only minor changes to the initial and the optimal structures of electrode compounds. To this extent, wellstructured and porous or layered nanomaterials are particularly interesting because of the shorter diffusion pathways they provide and the voids within the particles that are advantageous for compensating volume changes during the charge and discharge processes [3, 4]. Ion batteries, which are known for their reusability and cost-effectiveness, generally utilise cations as charge carriers. In this regard, the lithium-ion battery (LIB) is the key technology, and has shown ubiquitous utility, similar to its widespread integration in clean energy grids, in cellular phones, and laptops. The use of lithium in ion batteries is particularly interesting because it is the lightest metal with the largest standard potential. In addition, its small ionic size confers high mobility and diffusion rate. However, LIB expansion has been limited, owing to crucial drawbacks, including rapid ageing with time or charge/discharge cycles and high manufacturing costs [5]. New candidate materials for cathodes and anodes have been extensively researched in the past few years. The overall aim is to achieve high energy, high power densities, and fast charging, as well as to enable more efficient applications of renewable energy sources [6]. However, a fundamental understanding of the intrinsic limitations of most electrode materials remains scarce because of the complexity of monitoring the phenomena occurring during battery operations over a wide temporal range using the so-called operando characterisation techniques [7, 8]. While discharge and charge processes occur in millisecond fractions on atomic level, degradation phenomena, contributing to the capacity decay, span from fractions of seconds to months or years. On the other hand, chemical reactions at the interfaces occur simultaneously in fractions of nanometres to the bulk range. Mechanical degradation, caused by volume changes in electrode materials during redox reactions, poses a severe issue. With such variations in surface area, the behaviour of electrode materials varies in relation to their reactivity with electrolyte solutions. This illustrates the complexity associated with processes in battery materials and the necessity to investigate transformations occurring at various lengths and timescales [7, 8]. To monitor the impact of such operational structural changes, electrode materials can be removed from the battery cell and characterised ex situ. However, this approach can expose them to air and subsequent possible damages such as chemical alteration or structural changes, which render observation of the internal processes difficult. In the past few decades, different analytical techniques have emerged as highly promising methods that allow the non-destructive characterisation of battery materials, including

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chemical, electronic, and morphological parameters, under operando conditions [3]. For this purpose, synchrotron X-ray sources are especially advantageous because of their high power and flux, which allow X-rays to penetrate battery cells during their operation [3, 7]. The importance of operando studies on rechargeable batteries is especially reflected by the increase in the number of scientific publications focusing on these studies in the past several years (Figure 7.1). Moreover, the challenge of optimising ion batteries has been addressed by the search for alternatives to well-established LIBs. Such efforts have focused on the utilisation of more abundant alkali metals such as sodium and potassium, or other multivalent metal ions such as magnesium or zinc as charge carriers. For example, sodiumion batteries (SIBs) require relatively low manufacturing costs, resulting from the higher abundance and affordability of sodium, compared to lithium. However, the use of SIBs instead of LIBs faces several drawbacks, especially high redox potential, higher atomic mass, and larger ionic radius of sodium, compared to lithium, which result in a decrease in the theoretical energy density. While these disadvantages lessen the competitiveness of SIBs for most mobile applications, they remain of considerable interest for stationary power grids [9]. In summary, this chapter offers an overview of the different techniques available for studying the events that occur during the battery charge or discharge processes. These include chemical reactions and phase transformations as well as their influence on the electrochemical properties and performance of the main types of commercial and most-promising ion batteries.

Figure 7.1: Number of publications reporting operando studies on battery materials up to 2021 (data source: https://scifinder.cas.org).

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7.2 Theoretical background: intercalation and conversion reaction batteries 7.2.1 Intercalation reaction mechanisms Classic LIBs contain, among others, three main components, namely, a positive electrode (cathode), a negative electrode (anode) in which the Li+ ions intercalate (Figure 7.2a), and a lithium-containing electrolyte. In addition, microporous separators are generally included between the electrodes and electrolyte to prevent short circuits and simultaneously allow the permeability of Li+ ions [3, 10]. Cathode materials are often composed, for instance, of transition metal oxides, which can store lithium within a reversible guest–host insertion in an ideally well-defined and stable crystal structure. The oxidation of the transition metal ions results in the removal of Li atoms during charge. In general, the existing applied cathode materials in LIBs are classified into two main classes: i) anion close-packed or almost close-packed groups and ii) open-structure groups. In the first group, the redox active metal ion occupies sites between the anion sheets alternatively, offering the advantage of high energy stored by volume ratio units. The second group of substances includes more open structures such as vanadium oxides or “tunnel-like” compounds such as manganese dioxides and lithium iron phosphates [5]. Graphite is the most commonly used anode material in LIBs and operates via intercalation mechanism. In addition to its high electrical conductivity, low cost, and eco-friendly nature, this material offers an extremely high cycle life and exhibits important volumetric, specific, and practical capacities, in addition to the highest speed of electron transfer among most carbonaceous compounds [11, 12]. In detail, during battery charge, Li+ ions are extracted from the cathode (e.g. LiMO2) and transported into the anode (e.g. graphite C6) through the electrolyte: LiMO2ðsÞ ⇄ Li1−x MO2ðsÞ + xLi + + xe−

(7:1)

xLi ++ xe− + xC6ðsÞ ⇄ xLiC6ðsÞ

(7:2)

The charge process is thermodynamically disfavoured, requiring the application of a high potential. The discharge process is the reverse reaction, i.e. the deintercalation of Li+ ions from the anode and intercalation in the cathode. Notably, electrons do not pass through the electrolyte but are compensated by an external circuit, with electrical current flowing for both reactions. Moreover, during the first charge and discharge cycles, organic compounds within the electrolyte can be decomposed to form insoluble materials at the interface between the anode and electrolyte, forming a passivating layer known as the solid electrolyte interphase (SEI) [14]. For optimal performance, small amounts of conductive additives can be added to protect the cathodes from dissolution and overcharge and to facilitate the formation

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and stability of the SEI. In addition, an aluminium current collector usually backs the cathode, whereas a copper current collector usually backs the anode to supplement the electrodes, which possess moderate electrical conductivity, to supply electrons to the external circuit [3]. However, improvement of both the cathode and anode materials is required for building high-capacity batteries. Despite the complexity of the consequent electrochemical decomposition of organic salts in electrolytic solutions, high-potential cathodes with voltages up to 5 V, such as layered and spinel-type polyanion oxides, have been widely developed as insertion materials [15]. Significant efforts have also been made to improve the anode performance, which is limited by self-discharge caused by corrosion and by a comparably low specific capacity of graphite (372 mAh/g) [16]. Higher specific capacities can be obtained, exploiting alternative charge storage mechanisms such as conversion reactions [17], as explained below.

7.2.2 Conversion reaction mechanism During discharge of conversion-type cathode materials such as nanosized transition metal chalcogenides, the transition metal cations are reduced to their elemental state while the charge carriers such as Li+ ions form at least one inorganic compound with the anions (Figure 7.2b). In an ideal conversion, this process is reversible, and the nanosized parent compound as well as the initial transition metal oxidation states are reobtained: nLi+ + ne− + Mn + Xm ⇄ M0 + nLiXm=n

(7:3)

(X = O, S, Se, F, Cl, N, H, or others) This phenomenon is usually associated with phase transformations and further amorphisation of electrode materials. In a few compounds, an intercalation reaction occurs before conversion [18]. In general, conversion-type electrodes possess high theoretical gravimetric capacities of approximately 500–1,500 mAh/g [19]. The disadvantages of conversion materials include large volume expansion, sluggish kinetics, which yield poorly-resolved voltage plateaus, and poor Coulombic energy efficiencies. If these difficulties can be overcome, several transition metal compounds are attractive electrode materials [19]. For example, binary metal fluorides offer high specific densities, and are electroactive above 3 V versus Li+/Li; thus, these represent good candidates as cathode materials. This far, most conversion-type electrodes have been used as anode materials because of their rather low redox potentials [3]. To fully understand the large differences in performance among the various electrode materials with similar compositions and the mechanisms by which these materials function and fail, no matter if it is intercalation, conversion, or alloying, real-time monitoring of the structure of these compounds during several cycles of charge and discharge processes with the aid of operando characterisation techniques is required [24].

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7.3 Operando X-ray diffraction X-ray diffraction (XRD) is a powerful analysis tool that can be used for the investigation of a wide range of materials; this tool is particularly suitable for crystalline compounds. The information provided by X-ray techniques ranges from material identification over quantification to structural elucidation, including determination of lattice parameters, contents of different crystalline phases in a mixture, and microstructure information (domain size, strain, and texture). A powerful source of X-rays, such as synchrotron light, offers high brilliance and collimation, as well as wide tuneable energy. Penetration within the battery cell allows the acquisition of additional information on electrode materials, with a high spatial and temporal resolution. These features allow the study of a wide range of samples and suitable measures for commercial batteries, such as coin cells, without modification [3]. Operando XRD experiments can also be performed using high-intensity laboratory diffractometers equipped with a Mo- or Ag-Kα source, as demonstrated by Bartsch et al. [25] and explained in detail below.

a)

b)

Figure 7.2: Schematic representation of the reaction mechanism occurring in the charge and discharge process within (a) insertion-type and (b) conversion-type battery materials (reprinted from reference [13] by permission of Springer Nature, Copyright 2008).

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In general, operando XRD measurements are especially important because the structural changes on promising materials for battery electrodes must be controlled to avoid safety issues and to ensure a long cycle life. Under abusive real-world battery conditions, electrode materials may undergo undesired phase transformations that result in elevated operational temperatures, overcharging, or discharge. In particular, high temperatures are critical for phase transformations involving loss of oxygen and risks of catching fire or exploding. Reversible chemical processes can lead to capacity loss, accompanying changes in lattice parameters or unit cell volumes, thereby creating elevated strain or particle fracture. Decomposition of composite electrodes can also occur or risk-bearing metal plating during deep discharge, thus determining the voltage limits particularly important for ion batteries [3].

7.3.1 Lithium-ion batteries (LIBs) Owing to their numerous advantages, LIBs are the most commercialised ion batteries and have various applications, for example, in the electromobility sector. Fast charging, high energy, and high power densities are enabled by LIB technology, which derives 0 from the fact that lithium is the most electropositive (ELi + =Li = −3.04 V vs. standard hydrogen electrode, SHE) and the lightest metal [6]. LiCoO2 (LCO) was the most commonly encountered cathode material in LIBs, unchallenged for decades since its first commercialisation by Sony in 1991, because it exhibits a decent theoretical specific capacity of 274 mAh/g, high operating voltage, and good cycling performance [26]. However, detrimental phase transformations after delithiation of more than ≈ 0.5 Li/LCO limit the practical specific capacity to ≈ 140 mAh/g [26]. Besides, the main limitations of LCO are the high cost, toxicity, and raw material criticality of the constituent cobalt, as well as low thermal stability caused by the release of oxygen at high temperatures, heating, and the resulting safety issues. Alternatively, LiNiO2 exhibits the same crystal structure as LiCoO2, similar theoretical specific capacity, higher energy density, and lower cost, in comparison to those of LCO. However, LiNiO2 is less stable upon oxidation. By contrast, LiMnO2 combines thermal stability and cost-effectiveness, thus showing potential as a cathode material [27]. The problem with the latter is the very low cycle stability of Mn-combining electrodes (Mn leaching) [28, 29]. Interestingly, Ni, Mn, and Co can be combined with lithium in ternary metal oxides, with the formula LiNixCoyMnzO2 (x + y + z = 1, NCM). Depending on the chemical composition, NCM cathodes offer high volumetric energy densities and good thermal stability at a lower cost, compared to LCO [9, 20, 30]. However, charging of conventional NCMs to higher voltages leads to accelerated capacity fading and poor cycling abilities, depending on the chemical composition, thereby justifying the development of other derivatives of NCM materials [30]. Lithium- and manganese-rich NCMs (LMR-NCMs), which have been described either as solid solutions (i.e. Li1+xM1–xO2, where M = Ni, Co, and Mn) [21] or as composites of Li2MnO3 (x(Li2MnO3) · 1 − x(LiMO2))

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[22]. LMR-NCMs exhibit high specific capacities, and its cycling stability depends on the value x and the relative amounts of Mn, Ni, and Co in LiMO2 [23]. In addition to NCMs, various other intercalation-type materials have replaced the LCO cathode, including LiNixCoyAlzO2 (x + y + z = 1, NCA), LiMn2O4 (LMO), and LiFePO4 (LFP). Particularly, the latter cathode material regained enormous interest for applications in the low-cost battery sector due to the long cycle life and raw material abundance [31]. Operando XRD is an excellent tool to investigate the intercalation-type battery cells because the structural integrity of the parent compound is mainly preserved during cycling. One example is reported by Schweidler et al. [31] showcasing operando XRD measurement during long-term cycling of a LIB cell, consisting of Li1+x(Ni0.85Co0.1Mn0.05)1 −xO2 (NCM851005) as the cathode and graphite as the anode material. Although this Ni-rich NCM material enables enhanced specific capacities for the battery, the increased Ni fraction leads to Ni oxidisation during charging, generating highly reactive Ni4+ ions. This species, consequently, results in electrolyte decomposition on the NCM particle surface and gas evolution, leading to increased cell pressure; hence, the cell performance deteriorates during cycling. In addition, particles with high Ni concentrations suffer more critical volume changes during de-/lithiation, resulting in interfacial fracture and mechanical stress during cycling [31]. Operando XRD results (Figure 7.3) demonstrate that the interlayer distances reversibly increase and decrease along the c-axis (003 reflection) during the first 100 charge and discharge processes, evidencing the successful de-/intercalation of Li+ ions from/into NCM851005. After 150 cycles, the 003 reflection becomes broader, indicating the formation of two different NCM phases, which were identified as the hexagonal phases, H2 and H3, by Rietveld analysis [32]. The mechanism fails with the ongoing cycle number; thus, after 500 cycles, the initial reactions are no longer applicable. Moreover, the 002 graphite reflection was monitored (Figure 7.3b); from this information, the interlayer spacing of the anode material can be estimated, thereby providing information on the degree of de-/lithiation [31]. Ongoing research on LIB electrode materials is also aimed towards lightweight lithium–sulphur (Li–S) batteries, which are now renowned for their high specific energy (≈ 2,600 Wh/kg), low operating temperatures (reaching − 50 °C), and cost-effectiveness, owing to high sulphur abundance [3, 33]. In this type of battery, sulphur S8 or carbon/S8 composites are used as the cathode material, whereas Li metal is used as the anode material, producing a lightweight and low-cost system. During the discharge process, Li undergoes oxidation to form Li+ ions, which diffuse to the cathode for charge compensation, and a series of transformations occur until insoluble Li2S is formed as the final product at the cathode. During charge, Li2S is dissolved and Li metal is re-deposited on the anode. The overall reversible chemical reaction (Equation 4) corresponds to the formation of Li2S from the reaction of lithium and sulphur during discharge and the reverse reaction during charge [34]:

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150 75 0 150 75 0 150 75 0 150 75 0 150 75 0 150 75 0 150 75 0 150 75 0 8.2

b) 1st 100th 150th 200th 250th 300th 500th 750th

8.4

8.6 8.8 9.0 2 Theta / °, Mo-Kα1,2

9.2

Intensity / counts

Intensity / counts

a)

60 40 20 60 40 20 60 40 20 60 40 20 60 40 20 60 40 20 60 40 20 60 40 20 10.5

LiC6

LiC12

C

1st 100th 150th 200th 250th 300th 500th 750th

11.0

11.5 12.0 12.5 2 Theta / °, Mo-Kα1,2

13.0

Figure 7.3: Changes of (a) the 003 reflection for NCM851005 cathode and (b) the 002 reflection for the graphite anode during cycling. The blue colour represents the measurements recorded before starting the charging process under OCV, while the light blue colour shows the measurements in the charged state. The dark blue colour indicates measurements in the discharged state (reprinted with permission from reference [31]. Copyright 2019 American Chemical Society).

16Li + S8 ⇆ 8Li2 S

(7:4)

However, the reactions occurring within the cells for both processes are considerably more complex. They also involve the formation of polysulphide species, Li2Sx (8 ≤ x ≤ 4), which is hardly detectable because of their solubility and mobility in electrolytes. The aforementioned polysulphides are also responsible for anodic aging as they diffuse towards the Li metal anode; they drive corrosion and account for the loss of the active material. According to operando XRD analysis, Cui et al. [35] investigated the evolution of crystalline phases within a Li–S battery cell designed as described above. Analysis of the diffraction patterns shows the complete disappearance of characteristic reflections for elemental sulphur during discharge, indicating the complete reduction of the anode material and formation of lithium sulphide. Surprisingly, while previous findings in the literature suggest that most of the lithium sulphides remain insoluble during the subsequent charge [36–38], the authors reported the reappearance of reflections, corresponding to sulphur, indicating that the recrystallisation of sulphur was not altered by the rate of the galvanostatic electrochemical cycle (Figure 7.4). These notable results were compared to findings obtained using another synchrotron technique, operando transmission X-ray microscopy, which displayed subtle changes in the morphology of the

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electrode particles. The noticeable increase in the porosity of the material indicated that few polysulphides were formed. This potentially causes the deterioration of cell performance over several discharge–charge cycles [35].

Figure 7.4: Operando XRD patterns of a Li–S cell, cycled at C/8 and with a sulphur/Super P composite prepared as cathode: (a) XRD pattern at the start of discharge; (b) XRD patterns for the region of the Q-space, marked by the red box in a) for points a–j labelled in (c) the corresponding electrochemical plot. The XRD patterns show the reappearance of sulphur reflections at the end of the first full cycle. Sulphur reflections in (b) are labelled with their Miller indices. Unlabelled reflections originate from the pouch cells polyester and separator. XRD patterns in blue include sulphur reflections. The total discharge capacity is 755 mAh/g and the total charge capacity is 707 mAh/g (reprinted with permission from reference [35]. Copyright 2012 American Chemical Society).

Despite the advantages inherent to Li–S batteries, they suffer from short cycle life, precipitation of electronically insulating Li2S, which prevents the electrical current from passing through the electrodes, and high self-discharging rates, owing to the mobility of soluble polysulphide intermediates. In general, operando XRD analysis offers a deep understanding of the formation and dissolution of crystalline side products, and provides meaningful insights into the capacity fade of such Li–S batteries. Improvement efforts can be additionally made to address certain issues: for example, the use of additives in the electrolytic solution to suppress the mobility of polysulphides, and installation of polymeric or ceramic membranes to protect the Li metal anode and prevent crossover

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processes. The effectiveness of this strategy has been demonstrated by operando XRD analysis of a Li–S battery cell with crystalline α-S8 as the cathode and a silica-made glass fibre separator [39]. In this case, the results show that at the beginning of the experiment, only sharp reflections of α-S8 can be observed (red curve in Figure 7.5a). However, these reflections quickly disappear at the open-circuit voltage (OCV) as the battery is discharged, and lithiation occurs. During the cycle, two broad reflections appear at 25.6° and 28.3° 2θ (dashed lines in Figure 7.5c), which are attributed to silica-adsorbed polysulphides, referred to as PS1 and PS2. These reflections are still visible at t = 50 h, which corresponds to the end of lithiation (discharge), in addition to one broad reflection around 27° 2θ, accounting for Li2S (dashed-dotted line in Figure 7.5c). The cycle proceeds

Figure 7.5: Operando XRD measurements of a Li–S cell. (a) Waterfall representation of XRD patterns and (b) the corresponding galvanostatic curve, which was recorded during the first cycle of the Li–S cell at C/50 rate. The coloured patterns in (a) indicate major changes occurring during cycling of the Li–S cell and correspond to the coloured dots in (b), representing the different states of charge. (c) XRD contour plot of the data shown in (a), with the same galvanostatic curve as shown in (b). The intensity chart is shown to the right of (c). The asterisk refers to a reflection arising from a cell component. α-S8 and β-S8 are represented by white horizontal lines with diamond and oval symbols, respectively. The positions of the peaks, labelled PS1 and PS2, are indicated by vertical black dashed lines. Li2S refers to the solid end-oflithiation product, lithium sulphide, and is symbolised by a dashed-dotted black vertical line. Overall, the horizontal dotted black lines are used as a guideline for recording the changes in peak intensity as a function of potential/time (reprinted by permission from reference [39], Copyright 2017 Nature).

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with charging from t = 50 h to 90 h, and reflections of Li2S decrease until they completely vanish after t = 70 h. At the end of the charge, recrystallisation of another phase occurs, which is assigned as β-S8 phase [39]. To explain the formation of PS1 and PS2, the adsorption of polysulphides onto the silica surface of the glass fibre separator was suggested, and further confirmed by the design of an electrochemical cell, where silica was used as an additive in the electrolyte solution. Moreover, the scavenging of lithium polysulphides could effectively increase the specific charge density by up to 25%, compared to that of a similar Li–S cell, without the addition of silica to the electrolyte. Although the use of silica as a scavenging agent for polysulphides does not completely prevent the loss of cathode active materials, this approach provides a simple solution to the corrosion of the lithium anode and a promising pathway for further investigations in mitigating the drawbacks currently encountered for Li–S batteries [39].

7.3.2 Sodium-ion batteries (SIBs) As mentioned above, LIBs dominate the market for portable electronics, including cell phones, laptop computers, and electric vehicles, owing to their good performance for moderate to high energy demands for lightweight devices with fast charging ability. However, limitations of LIBs include safety issues, flammability of the organic electrolytes, and the presence of toxic components, making their disposal after the end of battery life a serious environmental concern [40]. Moreover, the market demand for lithium is constantly growing, but its reserves are neither renewable nor abundant. Given the scarcity of concentrated material sources, focusing the demand for large-scale applications on LIBs is impractical [41]. Alongside efforts to improve the properties of LIBs, research intensifies the focus on earth-abundant alternative compounds for safe, inexpensive, and environmentally friendly battery materials [40]. For lightweight materials, sodium-ion batteries (SIBs) are a possible alternative to LIBs. Importantly, there is an abundance of sodium in the Earth (2.38%) – more than 200 times higher than that of lithium (0.01%). Because sodium has also physical and electrochemical properties comparable to those of lithium, SIBs provide a possible solution for large-scale energy-storage systems [42]. For example, the redox potential 0 for sodium (ENa + =Na = −2.71 V vs. SHE) is only slightly higher than that of lithium 0 (ELi+ =Li = −3.05 V vs. SHE). However, this difference results in a loss of power and energy density [43]. Moreover, the development of SIBs has been hampered by the larger atomic radius of sodium compared to lithium, which crucially affects the electrochemical behaviour of structural distortions within the electrode materials [8, 40]. In general, layered NaxTMO2 (TM = transition metal or a combination of transition metals) are described in the literature as promising high-performance cathode materials for SIBs used in ambient temperature applications [41]. These sodium metal oxides adapt several polytypes,

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i.e. they differ in the stacking arrangement of the metal oxide layers, which depends on the transition metal TM and the sodium insertion amount x. In the P-polytypes, the Na+ ions are prismatically coordinated, whereas in O-polytypes, the Na+ ions are octahedrally coordinated. Particularly, in P2 structures, for example, common for NaxMnO2, the NaS6 prisms share rectangular faces, facilitating Na+ ion transport through wide passages with low activation energy (Figure 7.6, left). This drives rapid diffusion of Na+ ions, which can be advantageous for high current rate capabilities. The introduction of additional electrochemically active elements and the formation of bimetallic sodium transition metal oxides have been reported to avoid or delay irreversible phase transitions between P- and O-polytypes, thereby improving the cycle life of batteries [41]. To investigate the relationship between the current rate stability and the structural alterations, Goonetilleke et al. [41] performed operando XRD studies on Na2/3 Mn0.8Fe0.1Ti0.1O2 electrodes, which were cycled at different current densities vs. Na+/ Na in a voltage window of 1.5–4.2 V. The results (Figure 7.6, right) reveal the formation of a second P2-type phase, with smaller lattice parameters, herein called P2✶, when the battery cell reaches the charged state; eventually, this phase disappeared upon discharge. According to the authors, the presence of these two phases indicates the inhomogeneity of the electrode material, in which a change in the sodium concentration causes phase separation. Particles closer to the electrolyte interface react faster at high current densities than those closer to the current collector [41]. Operando XRD measurements during cycling of such layered NaxTMO2 compounds were also reported by Maletti et al. [44], who investigated NaNi0.5Ti0.5O2 as the cathode material for SIBs. Similarly, Ti ions are expected to suppress the transformation between the P- and O-polytypes. The results of these operando XRD experiments reveal reversible transformation between O- and P-type NaxNi0.5Ti0.5O2 during charge and discharge vs. Na+/Na. The structural changes during sodium extraction and insertion are described in detail, and the respective evolution of lattice parameters of both polytypes were tracked with respect to the variation of Na content x (Figure 7.7). The pristine O-type phase is observed for x = 1–0.65 and the P-type phase appears simultaneously for x < 0.85, whereas solely P-type phase is found for x < 0.6. During the coexistence of both polytypes, the lattice parameters of the P-type phase remain almost constant during Na insertion and extraction, while further Na is simultaneously inserted or extracted, respectively, from the O-type phase, which is reflected by the changes of lattice parameters along the a- and c-axes. The P-type phase starts to release or insert Na, only for x < 0.6 [44]. Lin et al. [8] performed operando XRD measurements during (de)sodiation processes on a nitrogen‐doped porous carbon-coated nickel cobalt bimetallic sulphide hollow nanocube (Ni0.5Co0.5)9S8@NC composite. As described above, many transition metal sulphides undergo conversion reactions, which is interesting for SIBs because of their enhanced reaction kinetics compared to oxides originating from relatively weak metal–sulphur bonds. The electrochemical performance of these materials can be enhanced by applying bimetallic transition metal sulphides instead of single-metal

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Figure 7.6: Structure of P2-type NaxTMO2 compounds, viewed along the a-axis (left), and formation of a new P2✶ phase, with smaller lattice parameter (higher 2θ values), during operando XRD measurements (right) (republished with permission of Royal Society of Chemistry from reference [41]; permission conveyed through Copyright Clearance Center, Inc.).

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Figure 7.7: Changes of lattice parameters a (top) and c (bottom) of NaxNi0.5Ti0.5O2 for the pristine O- (circles, R−3m) and P-polytype (squares, R3m) during charge (full symbols) and discharge (hollow symbols) of a SIB (reprinted with permission from reference [44]. Copyright 2019 American Chemical Society).

sulphides because of a richer redox chemistry and because more phases are formed during the first Li or Na uptake, yielding a more finely distributed nanostructure, with the conversion products being in close contact [44]. While reducing the transition metal sulphide particles to the nanoscale shortens the Na+ ion diffusion pathway, embedding the particles in heteroatom-doped porous carbon shells offers multiple advantages. On the one hand, the carbon shells increase the overall electronic conductivity of the composite and support the charge transfer to potentially insulating conversion intermediates and products, thus offering increased accessibility of active sites for charge storage. On the other hand, the carbon shells can adsorb polysulphide intermediates and prevent nanoparticle agglomeration. The reported (Ni0.5Co0.5)9S8@NC composite was prepared via Na2S sulphidation of phenol formaldehyde-coated Ni3[Co(CN)6]2 metal–organic framework precursors, followed by calcination [8]. Figure 7.8 shows the XRD patterns recorded in an operando measurement; the characteristic reflections of the (Ni0.5Co0.5)9S8@NC material are visible and reflections from the carbon cloth substrate are marked with asterisks. At the start of sodiation (between ≈ 2.5–0.75 V), the (Ni0.5Co0.5)9S8@NC phase seems stable, as corresponding reflections show no shift and show only a slight intensity decrease. During further sodiation (0.75–0.30 V), the intensity of the reflections of (Ni0.5Co0.5)9S8@NC strongly decreases and they finally disappear. Simultaneously, a new reflection at 39.0° 2θ evolves, which is assigned to the (220) lattice plane of Na2S, followed by increasing intensities for reflections at ≈ 45.5° and ≈ 54° 2θ, which indicate formation of Ni and Co metals. A set of reflections at 31.3°, 53.5°, and 56.3° 2θ was attributed to Na2S5 [8]. This study also shows that the resulphidation of metals and the re-formation of the

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(Ni0.5Co0.5)9S8@NC composite are incomplete, resulting in the formation of an amorphous material fraction. Additional investigations for SIBs should address the stability of the electrode materials during charge or discharge, and the improvement of the cell performance could be further elucidated by complementary in situ/operando experiments to provide a better understanding of molecular and structural dynamics, and possible pathways for the optimisation of these highly promising alternative batteries [8].

7.3.3 Potassium-ion batteries (PIBs) In the family of alkali metals, potassium-ion batteries (PIBs) have recently been explored as alternatives to LIBs as well. One important advantage of PIBs is the low reduction potential of potassium redox active species in non-aqueous electrolytes, 0 compared to that of Li (EK0 + =K = −2.93 V vs. SHE and ELi + =Li = −3.05 V vs. SHE). Additional PIB advantages include potentially low production costs, high abundance of potassium in the Earth’s crust, and easy insertion of K+ ions in graphitic anode materials [45]. The most commonly reported electrode materials for PIBs are layered metal oxides because of their stability, ready availability at low cost, and high theoretical energy density [46–48]. However, layered metal oxides suffer from low operating voltage and short cycle life [49]. Hexacyanometalates, AxMy[M(CN)6]z · nH2O (A = alkali cations, Zn2+, Al3+; M, M = transition metal ions), also known as Prussian Blue analogues (PBA), have shown promise for application in SIBs and PIBs because they combine long cycle life and high energy efficiency, either in aqueous or non-aqueous electrolytes during battery operation [50, 51]. The structure of PBAs consists of a rigid three-dimensional network, formed by an octahedron-based polyhedron around the metal cations. The voids formed within this structure are filled with zeolitic water molecules or alkali ions, and K+ ions can be electrochemically inserted into and extracted from the voids (Fig. 7.9a) [52]. These compounds can be easily synthesised by co-precipitation without calcination. The synthesis requires simple metal salts as precursors, and their variety allows tailoring of the electrode performance by selecting a suitable type of metal ion or by controlling stoichiometric ratios [53]. However, the viability of such PIBs depends on the choice of the electrolyte to address the problem of low Coulombic efficiency, upon degradation of the electrolyte at high potentials [54]. To optimise the performance of a potassium manganese hexacyanoferrate K2Mn[Fe (CN)6] as a cathode material in a PIB electrochemical cell, Pasta et al. [52] used an intrinsically stable electrolyte solution, obtained by dissolving potassium bis(fluorosulfonyl) imide (KFSI) in the ionic liquid N-butyl-N-methylpyrrolidinium bis(fluorosulfonyl)imide (Pyr1,3FSI). The intercalation of K+ ions in the anodic graphite host structure was probed by operando XRD during the first galvanostatic cycle. During discharge, the potential rapidly decreases from 2 to 0.33 V, and subsequently two pronounced pseudo-plateaus

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are observed during discharge to 0.0 V. As seen in Figure 7.9, these two pseudo-plateaus can be attributed to, firstly, the formation of three intermediate phases (KC48, KC36, and KC24), and secondly, the generation of the most K-rich fully phase KC8, as the reflections of the pristine graphite evolve towards a diffraction angle of 33.5° 2θ. This transformation is fully reversible and the same intermediate phases appear during K removal, and finally, the recovery of solely graphite [52].

Figure 7.8: (a) Schematic representation of the synthetic steps for the preparation of the (Ni0.5Co0.5)9S8@NC composite, including coating Ni3[Co(CN)6]2 with phenol formaldehyde (RF), sulphidation with Na2S, and final calcination at 550 °C. (b) Operando XRD results during the first discharge and charge cycle for (Ni0.5Co0.5)9S8@NC-based electrodes (reproduced with permission of Wiley from reference [8], permission conveyed by the Copyright Clearance Center Inc.).

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Figure 7.9: (a) A schematic representation of the full-cell potassium-ion battery consisting of a graphite anode, potassium bis(fluorosulfonyl)imide in N-butyl-N-methylpyrrolidinium bis(fluorosulfonyl)imide electrolyte, and K2Mn[Fe(CN)6] cathode. (b) Operando XRD patterns and (c) corresponding potential vs. capacity profile for potassium intercalation and deintercalation into graphite at C/30. The major reflections for the intermediates KC48, KC36, KC24 and KC8 are highlighted (adapted with permission from reference [52]. Copyright 2020 American Chemical Society).

7.3.4 (Hybrid) zinc-ion batteries (ZIBs) Zn-based electrode materials are interesting because of their relatively low potential 0 (EZn 2 +=Zn = − 0.76 V vs. SHE), high theoretical capacity (820 mA h/g for Zn metal), and natural abundance of Zn [55]. One of the foremost advantages is the operation of Znbased batteries in aqueous electrolytes, drastically reducing safety concerns, compared to the flammable organic electrolytes necessary for LIBs, SIBs, and PIBs. Moreover, aqueous rechargeable batteries are attractive, owing to their low cost, environmentfriendly characteristics, and high conductivity of the electrolyte – crucial for large-scale energy storage [40, 55].

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However, a major challenge in the development of Zn-based batteries is the slow diffusion of Zn2+ ions, caused by its high charge density, thereby leading to strong Coulomb interactions within the polar crystalline hosts. Also, the formation of Zn dendrites is problematic, because they cause capacity loss and poor Coulombic efficiency, and deterioration of cycle life. The lack of information on the working mechanism behind the de-/intercalation of Zn2+ ions is another issue. Hybrid Zn-ion batteries (ZIBs) aim to overcome this difficult shift from single-ion to double-ion charge and discharge mechanisms [40, 55]. Guo et al. [40] investigated the working mechanism of an aqueous hybrid LFP/Znbattery by operando XRD, a schematic representation of the cell is shown in Figure 7.10a. Synchrotron radiation helped for this operando experiment to reveal that only Li+ ions participate in the de-/intercalation processes from/in the cathodic iron phosphate, while Zn2+ ions are deposited in the Zn anode, resulting in the following proposed working mechanism [56]: Cathode: LiFePO4 ⇆ e − + Li + + FePO4 Anode:

Zn2 + + 2e − ⇆ Zn

(7:5)

Overall: 2LiFePO4 + Zn2 +⇆ 2Li + + 2FePO4 + Zn In the operando XRD measurements (Figure 7.10b), the reflections assigned to LiFePO4 are detected in the pristine state, and disappear during charge. Simultaneously, reflections assigned to FePO4 start to appear, demonstrating full delithiation from LiFePO4 to form FePO4. During discharge of the cell, LiFePO4 is reobtained, which confirmed the successful uptake of only Li+ ions. Finally, sodium dodecyl benzene sulphonate was added to improve the mobility of Li+ ions and inhibit the growth of Zn dendrites [40]. In an additional study on Zn-based aqueous rechargeable batteries, operando XRD experiments were performed using K+ ion intercalated V2O5 (KVO) nanorods as cathode materials. Here, alkali metals such as potassium function as pillars within the structure of transition metal oxides to increase their interlayer spacing (Figure 7.11a). This strategy enhances the diffusion of Zn2+ ions within the ZIB and prevents the collapse of the layered cathode structure during guest insertion and extraction, thereby improving the battery cycle life. Shifts of reflections detected during the operando XRD measurements show that Zn2+ ions diffused along the a-axis of the KVO parent structure (Figure 7.11b). Because the (002) and (004) planes of the KVO nanorods were exposed to the electrolyte, they provided access for the intercalation of Zn2+ ions, enabling high discharge capacity and high-rate capability. Because of the occupation by K+ ions, the intercalation of Zn2+ ions along the b- and c-axes is improbable [55].

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Figure 7.10: (a) Schematic representation of the working principle of the hybrid LFP/Zn battery; (b) operando XRD measurement of this battery, with respective charge-discharge curve; and (c) contour plots (reproduced with permission of Wiley from reference [40], permission conveyed by the Copyright Clearance Center Inc.).

Figure 7.11: (a) Schematic representation of the storage mechanism of the KVO in the aqueous rechargeable ZIB and (b) detailed models of the KVO structure described for the (100) surface (republished with permission of Royal Society of Chemistry from reference [55]; permission conveyed through Copyright Clearance Center, Inc.).

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7.3.5 Magnesium-ion batteries (MIBs) In addition to batteries with lithium, sodium, potassium, and zinc ions, magnesium-ion batteries (MIBs) have been recently investigated in situ during operation. Advantages such as high abundance, low cost, high theoretical capacity, and high safety justify a possible LIB replacement for large-scale energy storage units, such as power grids [57]. Thus, Fu et al. [57] studied the structural evolution of orthorhombic (α-)V2O5 nanowires during the electrochemical insertion and extraction of magnesium. The α-V2O5 sheetlike structure is advantageous for intercalation reactions, and the reduction of the material dimension to the nanoscale is expected to reduce the diffusion barrier for Mg2+ ions. Therefore, possible phase transitions of V2O5 between the α, ε, γ, and ω polymorphs during cycling were investigated in detail. Operando XRD results (Figure 7.12) show that at the beginning of the first discharge vs. Mg2+/Mg up to a certain stoichiometry of Mg0.14V2O5, Mg2+ ions are incorporated into the V2O5 lattice through a solid solution mechanism. In this reaction step, an increase in the lattice parameters a and c and a decrease for b are observed. A two-phase mechanism follows, in which the fraction of Mg0.14V2O5 decreases while the fraction of Mg0.6V2O5 increases. During this period, the lattice parameters remain constant, yielding a phase ratio of Mg0.14V2O5: Mg0.6V2O5 of 13:87 at the end of the first discharge. During the first charge, the lattice parameters, a and c, decrease, while b increases because of a combined two-phase transition and solid-solution mechanism. A comparable picture is observed in the second cycle; however, a phase ratio of Mg0.14V2O5:Mg0.6V2O5 of 15:85 near the end of the second charge is observed, and the initial V2O5 parent structure is reobtained at the end of the second charge. In contrast to the results in previous reports, the ε- and γ-phases were not observed during cycling [57].

7.3.6 All-solid-state batteries (ASSBs) All-solid-state batteries (ASSBs) are highly interesting in the field of electric mobility, owing to their potentially improved safety and wider operating temperature range than those of conventional LIBs, as well as its potential to deliver very higher energy and power densities by implementation of Li metal anodes or concepts based on in situ Li metal deposition [25]. However, limitations such as instabilities at the cathode active material (CAM) and SEI, chemo-mechanical degradation, and development of scalable fabrication processes need to be overcome. One important example of this approach was reported by Bartsch et al. [25], who explored the combination of lithium thiophosphate solid electrolytes (SEs) and Ni-rich layered oxides as high-energy cathode active materials. Here, the degree of de-/lithiation of Li1 +x(Ni0.6Co0.2Mn0.2)1 −xO2 (NCM622) CAM was studied via operando XRD during charge and discharge by correlations between Li content and changes in the lattice parameters. Operando experiments were completed by investigating the effect of a LiNbO3 protective coating to

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Figure 7.12: Changes of MgxV2O5 lattice parameters (middle), Mg0.14V2O5:Mg0.6V2O5 phase ratios (bottom) and galvanostatic discharge-charge curves (top) during cycling (reprinted with permission from reference [57]. Copyright 2019 American Chemical Society).

prevent direct contact between the CAM and the thiophosphate SE, which is expected to partially decompose during electrochemical cycling. According to the results, the domain length along the a-axis decreases during delithiation because of Ni oxidation and the consequent decrease in the ionic radius, which causes a contraction of the abplane. By contrast, the domain size along the c-axis (Figure 7.13) first increases and subsequently decreases after approximately x(Li) ≈ 0.5. The authors attribute the first increase to the electrostatic repulsion between the oxygen layers, while the subsequent decrease is justified by the charge transfer between the oxygen and nickel atoms at a high state of charge (SOC). Also, the LiNbO3-coated CAM in these ASSB cells can de-/intercalate more Li+ ions (x(Li) ≈ 0.62), compared to uncoated CAM (x(Li) ≈ 0.51), mainly because of the higher delithiation degree of the coated cells in the initial charge cycle and enhanced lithiation during discharge [25].

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7.4 Operando X-ray absorption spectroscopy X-ray absorption spectroscopy (XAS) is a powerful tool to study alterations of electronic states and the local environments around specific elements. Operando XAS experiments have been combined with operando XRD measurements to elucidate the lithium storage mechanism in a conversion-type ZnCo2O4 (ZCO) electrode material [58]. Such transition metal oxides are advantageous for increasing the specific capacity by two or three

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times, in comparison to those of commercial graphite electrodes for LIBs. Operando results reveal the formation of LiCo2O3, ZnO and CoO as intermediate phases during the first lithiation process. Upon delithiation, these phases are converted to CoO and ZnO, instead of the re-formation of the initial ZnCo2O4. Subsequently, alloying with Zn occurs when the electrode metal ions are reduced to their elemental state [58]. A composite anode, based on ZnCo2O4, loaded on a carbon cloth (ZCO/CC) is a flexible, binder-free material. The composite anode delivers high specific capacities even after several galvanostatic cycles (e.g. 701 mA h/g at the 60th cycle at a specific current density of 0.25 A/g) [58]. Permien et al. [4] showed that, in contrast to earlier findings in literature based on conventional XRD data, nanosized NiFeMnO4 consists of a mixture of NiO and a strained cubic spinel phase [4]. This discovery was enabled by a combination of synchrotron XRD and pair distribution function (PDF) that corrected inaccurate assumptions of stoichiometry and oxidation states of this important electrode material, which possessed high reversible capacity of ca. 840 mAh/g up to the 50th cycle. Results from operando XAS, in combination with operando XRD, reveals that an Li uptake per formula unit (fu) of ca. 0.8 Li/fu is possible without changing the electronic structure. Mn3+ and Fe3+ ions were simultaneously reduced by increasing Li/fu, and move from tetrahedral to empty octahedral sites, before they are finally converted to their elemental state after further Li uptake [4]. Li et al. [59] used operando X-ray absorption near-edge spectroscopy (XANES) to study the interfacial performance between a commercial Ni-rich LiNi0.8Co0.1Mn0.1O2 (NCM811) cathode and Li10GeP2S12 SE. Despite their high ionic conductivities (10–2–10–4 S cm–1), most promising ASSB with sulphidic SE still suffer from severe issues. Some of these problems include limited cycle life and rapid performance degradation, resulting from the unstable interface between the electrode and the electrolyte material, and from chemo-mechanical degradation. Under these specific conditions, ex situ mechanistic studies to understand and prevent these issues are disadvantageous. Thus, a modified coin cell (Figure 7.14), with an open window to allow the penetration of incident X-rays, was used in these operando XANES experiments. Within the Ni Kedge spectra, the signal at 8353 eV, which was assigned to the 1s→4p transition, initially shifted to a higher energy and then returned to a lower energy level. According to the authors, these shifts occur because of Ni2+/Ni3+ and Ni3+/Ni4+ electrochemical redox reactions in the NCM811 cathode material during the delithiation and lithiation processes. Considering the S K-edge feature of the Li10GeP2S12 SE at 2470.7 eV, a gradual shift to lower energies is observed during the first charge. In the high-voltage state (> 3.5 V), the additional shift to 2470.4 eV indicates the instability of the electrolyte. During discharge, the feature at 2470.4 eV returned to 2470.7 eV, and a new feature at 2472.5 eV appeared. Interestingly, this new signal was assigned to the S 1s to Li2S σ✶ transition, suggesting that the electrolyte material is first decomposed to Li2S instead of to other metal sulphides or polysulphides [59].

Chapter 7 Operando studies on the charge and discharge processes of battery materials

Figure 7.14: (a) Schematic representation of the operando XANES battery cell. (b) Operando Ni and S K-edge spectra with the respective charge (red arrows) and discharge (blue arrows) profiles of a LiNi0.8Co0.1Mn0.1O2–Li10GeP2S12 ASSB (adapted with permission from reference [59]. Copyright 2019 American Chemical Society).

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7.5 Operando imaging techniques Yu et al. [60] performed operando X-ray imaging experiments at the Cornell High Energy Synchrotron Source with 10 keV energy and a resolution of 2 µm to study the plating and stripping of Li dendrites in Li metal electrodes. Li metal is an interesting anode material for rechargeable batteries because of its high specific capacity (3,860 mAh/g) and low reaction voltage. However, one of its major disadvantages is the growth of dendrites during plating, which cause inhomogeneities across the electrode surface and the consequent failure of the material [60]. Also, the formation of dendrites is an inherent safety issue because dendrite growth can lead to short circuits, which eventually ignite the flammable organic electrolytes and risk explosion and fire hazards of such batteries. This study aims to offer a systematic investigation of Li electrodeposition under practical battery conditions. Figure 7.15 shows the influence of the concentration of LiPF6-based electrolyte (0.5 and 0.1 M) on the morphology evolution of plated Li. At lower concentrations, the amount of plated Li is less under the same charge. By contrast, in 0.5 M electrolyte, the plated Li possesses a moss-like structure, while in the 0.1 M electrolyte, dense Li clusters are formed, turning into “cactus”shaped structures upon growth. Additional experiments have demonstrated the formation of new Li plating morphologies, such as “volcanic Li”, which originates from the use of CsPF6 in this study as an additive to the LiPF6 electrolyte, and pits on the electrode surface during Li stripping [60]. Li dendrite growth in LIBs has also been studied recently under operando conditions by Song et al. [61]; hence, neutron imaging techniques are particularly efficient, owing to the high penetration depth of neutrons. Li deposition on the electrode surface during long-term cycling is typically inhomogeneous, forming moss- or needle-like microscale dendrites. These dendrites penetrate the separator, forming a direct current path between the positive and negative electrodes thus causing short circuits. The work of Song et al. [61] particularly focuses on the consequence of internal short circuits, and the authors developed a model based on operando neutron imaging studies, which is summarised in Figure 7.16a. Stage 1 is characterised by the deintercalation of Li+ ions from the LMO host during charge. Subsequently, these Li+ ions are transported through the electrolyte to the Li metal anode, where they are reduced and deposited on the Li metal substrate. In stage 2, inhomogeneous nucleation causes the formation of dendrites, which grow further and penetrate the glass fibre used as separator. A bridge is formed with the delithiated Li1−xMn2O4 area, causing a short circuit. Consequently, the battery cannot be charged at higher voltages. This result explains the decrease in voltage shown in Figure 7.16b. Then, the internal resistance around the contact areas decreases, resulting in electron transport from the dendritic Li to the Li1−xMn2O4 area. Interestingly, the Li+ ions intercalated into Li1−xMn2O4 are subsequently reduced by these electrons, forming the LiMn2O4 phase again. Dendritic Li is further consumed by the reaction with Li1−xMn2O4 to form LiMn2O4. In stage 3, the dendrites have shortened considerably so that they lose contact with LiMn2O4. The electron transport pathway is

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interrupted, self-discharge stops, and the battery can be charged again (back to stage 2). However, the battery short circuit is reached shortly afterwards, and the charge–selfdischarge process becomes the dominant mechanism for battery cycling [61].

Figure 7.15: Images of the SEI layer evolution of plated Li in 0.5 and 0.1 M electrolytes under practical battery conditions (adapted with permission from reference [60]. Copyright 2019 American Chemical Society).

An interesting study on battery operando XRD computed tomography was reported by Liu et al. [62] for quantification of non-uniformities, not only as a function of time but also as a function of position. Such information is very helpful to optimise battery performance because local limitations in ionic and electronic transport can lead to non-uniform energy storage reactions, especially in composite battery electrode compositions. These composite architectures usually consist of a mixture of the active phase, conductive carbon additives, binder, and electrolyte-accessible porous regions, where heterogeneous performance causes non-uniform degradation and capacity loss. In the work of Liu et al. [62], the topographical mapping of an LFP-based electrode during charge and discharge reveals accelerated reactions at the electrode faces in contact with the separator or current collector. Surprisingly, the maximum reaction rate experienced by the individual cathode particles in these regions is two to five times higher than the average control charge and discharge rates [62]. Wu et al. [63] used a transparent electrochemical cell to visualise the operation of zinc–bromide batteries during cycling. These zinc-bromide redox-flow batteries (ZBFBs), which possess a high theoretical energy density, offer an interesting solution for stationary energy storage technology, and are currently on path to commercialisation [63]. Regarding the charge storage mechanism, the total ZBFB reaction can be written as follows:

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charge

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Figure 7.16: (a) Schematic representation of mechanistic pathway of Li+ ions during battery charge, short circuit, and self-discharge. Stage 1: deintercalation of the Li+ ions from the LiMn2O4 host during charge. Stage 2: inhomogeneous nucleation causes dendrite formation, and bridging with the delithiated Li1-xMn2O4 area results in short circuit and self-discharge. Stage 3: dendritic Li is consumed in the battery cycle. The dendrites become short and lose contact with LiMn2O4. The electron transport pathway is interrupted, and self-discharge stops. (b) Charging curve during the operando experiment (adapted with permission from reference [61]. Copyright 2019 American Chemical Society).

However, the generated Br2 product is soluble in aqueous electrolytes, resulting in crosscontamination and diffusion to the Zn anode side. This contamination decreases the Coulombic efficiency and causes health and environmental problems due to the high vapour pressure of Br2. A possible solution to this problem is the sequestration of Br2 by complexing agents such as N-methylethyl-pyrrolidinium bromide (MEPBr) to form polybromides (MEP + Br2n +1–, n = 1–4). The operating principle of this solution has rarely been studied because of practical challenges. For instance, similar to polysulphides, polybromides can be modified by extraction from the reaction environment. In addition, polybromides are highly sensitive to X-ray radiation and the vacuum environments in electron microscopes, thereby limiting the available methods for characterisation under real operando conditions. Here, dark-field light microscopy (DFLM) is advantageous

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Figure 7.17: (a) Schematic representation and photographs of the cell and microscopic images for operando visualisation of polybromide formation on zinc–bromine flow batteries. (b) Nucleation and (c) merging of polybromide droplets. (d) Disappearance of polybromides during discharge. (e) Sketch of a droplet immobilized on the Pt surface during discharge; and (f) voltage profile during the first cycle (reproduced with permission from Wiley from reference [63]).

because this technique offers high contrast and sensitivity, in addition to its accessibility and non-disturbing characteristic. For the operando microscopy experiments, the battery cell shown in Figure 7.17a was constructed using a 100 nm Pt layer on a glass substrate as the working electrode, carbon paper as the counter electrode, and a solution containing

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Br–, MEP+ and Zn2+ ions as the electrolyte. Sporadically distributed polybromide products is generated under a small oxidising current (Figure 7.17). After 5 min, spherical microparticles of 2–3 µm size are formed on the Pt surface, growing to ≈ 10 µm at the end of the charge and merging into larger-sized microstructures in contact with other microparticles. During discharge, the reduction to Br– ions causes the reduction of liquid droplets, decreasing only in thickness according to the assumption of an anisotropic pattern, and become transparent before they disappear. In summary, operando microscopy results reveal that the Br2 charge products are dissolved in the aqueous medium. These droplets are formed upon deposition at the same position on the electrode surface and are consumed thereafter during discharge (see Figure 7.17b–d). Interestingly, polybromide formation does not depend on the substrate, but is strongly influenced by side reactions induced by overcharging [63].

7.6 Operando optical and vibrational spectroscopy Chen et al. [64] studied the formation of Li dendrites on an unstable SEI and the resulting low Coulombic efficiency of Li plating and stripping. Here, a strategy was developed to deposit a multifunctional antimony-based interphase on a Li metal anode by immersing it in an antimony triiodide–tetrahydrofuran (THF) solution (Fig. 7.18). The resulting lithiophilic layer was composed of gradients of amorphous antimony (Li3Sb, Sb, and SbOx) and lithium (LiI, LiOH, Li2CO3, and Li2O) compounds. This interphase exhibits high ionic and low electronic conductivity, as well as good electrochemical stability. The homogeneous electronic insulation of lithium compounds on this interphase prevents the formation of Li dendrites. Operando Raman measurements recorded through a sealed quartz window in the applied Li–S battery cell confirms that the formation of this antimony–lithium-based interphase reduces parasitic reactions with the liquid electrolyte and relieves the transport of polysulphides [64]. Ren et al. [65] used an interesting new approach to detect generation of water traces in LIBs under operando conditions. Water impurities are usually formed by side reactions in liquid electrolytes during electrochemical cycling, resulting in poor cycling performance and loss of active materials. This generation occurs due to several reasons. For example, capacity fading can be triggered by the reaction of water with lithium-foil electrodes or LiPF6 electrolytes. On the other hand, water can also destroy the SEI. This reaction passivates the electrode surface and prevents electrolyte degradation, consequently allowing the transport of Li+ ions and resisting electron transport. Finally, the reaction between water and the anode materials can generate the inflammable hydrogen gas, thereby increasing the internal pressure of the battery [65]. In the work of Ren et al., the [Tb2(DHBDC)3(DMF)4] (DHBDC = 2,5-dihydroxyterephthalic acid, DMF = N,Ndimethylformamide) coordination polymer’s turn-on fluorescence response towards moisture was used for detecting water formation during battery discharge.

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Figure 7.18: (a) Schematic representation of the formation of the lithiophilic antimony-based layer to prevent Li dendrite formation. (b) Photograph and (c) schematic configuration of the Li–S cell constructed for the operando Raman measurements (reprinted with permission from reference [64]. Copyright 2019 American Chemical Society).

The authors proposed that the intrinsic emission of the DHBDC ligand is selfquenched in the pristine coordination polymer structure because of the short interligand distances, which are below the critical radius for Coulombic energy transfer. Upon exposure to moisture, the water molecules preferentially coordinate the Tb3+ ions, causing structural decomposition and separation of the DHBDC molecules, and stopping the quenching of their luminescent properties. However, the absence of the 4f→4f electronic transitions of the Tb3+ ions within the fluorescence spectra was not explained. For operando measurements, the coordination polymer was used as an additive in the electrolyte solution (Figure 7.19). In situ luminescence measurements reveal a gradual increase in emission, attributed to the linker, indicating water formation during the first discharge process. In total, approximately 0.18% of water is generated in 2 mL of the electrolyte solution, corresponding to the formation of a SEI layer [65].

7.7 Other operando techniques The performances of batteries during operation have also been investigated using small-angle X-ray scattering (SAXS) or small-angle neutron scattering (SANS). For example, Li–S cells were probed by Risse et al. [33]; the authors investigated the formation of S8 and Li2S, which reportedly cause capacity fading [66], within the pores of the carbonaceous cathode composite. Here, SANS is an ideal method to tackle this question, delivering quantitative information regarding the internal surface, size, and shape of the pores. In addition, the method is complemented by the high penetration depth of the neutrons. The operando cell (Figure 7.20a) was constructed from chemically inert polyether ether ketone (PEEK), stainless steel current collectors, and exchangeable aluminium windows, which are transparent to neutrons. Figure 7.20b shows the expected

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Figure 7.19: (a) Schematic representation of the applied operando cell and (b) correlation of in situ emission spectra with the simultaneously measured specific capacity (Royal Society of Chemistry from reference [65]; permission conveyed through Copyright Clearance Center, Inc.).

SANS curves if the precipitation of Li2S and S8 is not correlated with micropore scattering. As demonstrated by the respective operando results, both products precipitated not in the pores but on the outer surface of the carbon fibres [66]. In addition, a different approach using a combination of operando liquid secondary ion mass spectrometry and molecular dynamics simulations has been explored to elucidate the mechanism behind the formation of SEIs on anode surfaces – one of the least understood component in LIBs [67]. The work of Zhou et al. [67] showcases the initial formation of an electric double layer, caused by self-assembly of the solvent molecules, and directed by Li+ ions and the electrode surface potential (Figure 7.21). The negatively charged electrode surface repels salt anions, resulting in the formation of a thin, dense, and inorganic inner layer that conducts Li+ ions and insulates electrons; this inner layer is accountable for the main functions of the SEI. After the formation of this inner layer, the electrolyte-permeable and organic-rich outer layers increase [67].

Chapter 7 Operando studies on the charge and discharge processes of battery materials

Figure 7.20: (a) Schematic of the cell used for small-angle neutron scattering (SANS) operando measurements on Li/S batteries. (b) Expected SANS curves for precipitation of Li2S and S8 outside the pores of the applied carbonaceous cathode material (left) and illustration of particle formation on the cathode surface (right) (reproduced with permission from reference [33]. Copyright 2019 American Chemical Society).

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f)

Figure 7.21: A schematic illustration of the formation of an SEI layer on the Cu anode surface with 1.0 M lithium bis(fluorosulfonyl)imide (LiFSI) in 1,2-dimethoxyethane (DME) as the electrolyte the SIMS depth profiles from the reference are not shown in this Figure 7.21 (reprinted by permission from Nature of reference [67]).

7.8 Conclusions and Outlook At present, LIBs are the most widely used rechargeable batteries in the electromobility market, and their applications range from low-power appliances such as cellular phones or laptop computers to high-power systems such as (hybrid) electric vehicles [2]. Increasing energy demand necessitates an increase in LIB performance, which is currently accompanied by high cost, security failures, and limited by too low energy densities. The exhaustion of lithium, which represents only 0.01% of the Earth’s crust, is predicted to be accelerated by the absence of complementary systems. A wide range of materials have been developed as alternative candidates that can improve the applications of LIBs [68]. In the studies presented in this chapter, several electrode materials were investigated for enabling the replacement of Li+ ions by more abundant metal ions such as Na+, K+, Zn2+, and Mg2+ ions, which can function as charge carriers. However, when these promising rechargeable battery candidates are applied to successfully overcome the disadvantages of LIBs, the increased ionic sizes of the charge carriers and the reduced reactivity give rise to issues concerning the stability of the electrode materials and the operating potential range. Therefore, it is important to gain detailed insights into the fundamental electrochemical reactions, structural, electronic, and morphological changes. The events during battery operation are solely assignable if in situ conditions are applied. The operando characterisation of electrode materials during charge and discharge provides a better understanding of their scope and limitations. For example, the operando techniques discussed in this chapter permit the application of bimetallic NaNi0.5Ti0.5O2 as a cathode

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material in SIBs without detrimental irreversible phase transformations [44]. For SIBs and PIBs, alkali metal hexacyanometalates have emerged as promising cathode materials, owing to their high electrochemical stabilities and open structures; however, the operation of, for example, PBA requires a suitable electrolytic solution that is stable at high operating voltages [52]. On the other hand, an example for the investigation of ASSBs is given by an operando XRD study, where the coating of the CAM stabilises the interface with a SE, and offers higher de-/lithitation degrees. Operando techniques other than XRD have also been presented as complementary analysis tools to obtain a better understanding of the ion battery function. Operando XAS, for example, in combination with operando XRD, offers the elucidation of more complex charge storage mechanisms such as the conversion of nanosized NiMnFeO4 as anode material for LIBs [4]. Additional combinatorial approaches are presented, which open up avenues for deeper insights into the mechanisms during charge carrier uptake and release. For example, initial XRD analysis showed the recrystallisation of S8 after several galvanostatic cycles in Li–S batteries, but SANS evidenced that Li2S precipitated on the outer surface of the carbon fibres in this system. In summary, the operando techniques described in this chapter are powerful tools to understand the fundamental electrochemical reactions and offer detailed insights into charge storage mechanisms for battery materials. This is outlined by several examples given, which provide a bigger picture and showcase the strength of operando investigations. Hence, these tools open up avenues for further optimisation of LIBs and can also provide valuable information to tackle more general challenges of ion batteries, which need to be understood as a next step towards alternative energy storage devices such as SIBs, PIBs, ZIBs, and MIBs.

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