Handbook of Sol-Gel Science and Technology [2 ed.] 9783319320991, 9783319321011

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Lisa Klein Mario Aparicio Andrei Jitianu Editors

Handbook of Sol-Gel Science and Technology Processing, Characterization and Applications Second Edition

Handbook of Sol-Gel Science and Technology

Lisa Klein • Mario Aparicio • Andrei Jitianu Editors

Handbook of Sol-Gel Science and Technology Processing, Characterization and Applications Second Edition

With 1736 Figures and 209 Tables

Editors Lisa Klein Materials Science and Engineering Rutgers University Piscataway, NJ, USA

Mario Aparicio Instituto de Cerámica y Vidrio Consejo Superior de Investigaciones Científicas (CSIC) Madrid, Spain

Andrei Jitianu Department of Chemistry Lehman College The City University of New York Bronx, NY, USA Chemistry and Biochemistry PhD Programs - The Graduate Center The City University of New York New York, NY, USA

ISBN 978-3-319-32099-1 ISBN 978-3-319-32101-1 (eBook) ISBN 978-3-319-32100-4 (print and electronic bundle) https://doi.org/10.1007/978-3-319-32101-1 Library of Congress Control Number: 2017961164 1st edition: # Kluwer Academic Publishers 2005 # Springer International Publishing AG, part of Springer Nature 2018 This work is subject to copyright. All rights are reserved by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, express or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publisher remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. Printed on acid-free paper This Springer imprint is published by The registered company Springer International Publishing AG part of Springer Nature. The registered company address is: Gewerbestrasse 11, 6330 Cham, Switzerland

Sumio Sakka, Professor Emeritus of Kyoto University, Hirakata, Osaka, Japan. All of us in the sol-gel science and technology community are indebted to Prof. Sakka. He was instrumental in organizing the Journal of Sol-Gel Science and Technology and the original Handbook Volumes. He provided encouragement to these Editors to expand this project and to turn these volumes into a living resource.

Preface to the Second Edition

It is a short 13 years since the publishing of the Handbook of Sol-Gel Science and Technology. During this period of rapid technological change, especially with regard to nanomaterials, the Handbook has remained an excellent source for the fundamentals, with plenty of motivation for new developments. Indeed, the Handbook has survived the test of time. In updating and expanding the Handbook, our goal is largely the same as it was when it was first edited by Professor Sumio Sakka. The original Handbook had a total of 85 chapters, divided into three volumes. Most of these chapters appear in the new Handbook, many of them updated by the original authors. What has changed since that time is that the scope and range of applications has increased. While several chapters have been added in the areas of processing and characterization, the number of chapters on applications has more than doubled. The applications are grouped in seven broad areas: mechanical, optical, electrical, electrochemical, preservation, organic-inorganic hybrids, and bio-related materials. The total number of chapters in this version is 125. With this updated and expanded Handbook, we now have the opportunity to make this a “living” Handbook, meaning that we can update and expand on a continuous basis. Nevertheless, we are staying true to the vision that Professor Sakka expressed so well in the Preface to the Handbook in 2005, included below. Mario Aparicio Andrei Jitianu Lisa Klein

vii

Preface to the First Edition

This three-volume Handbook “Sol-Gel Science and Technology” was planned with the purpose of providing those who are interested in processing, characterization and application of materials, with comprehensive knowledge on sol-gel science and technology. Around 1970, three different groups in the field of inorganic materials published research results on preparation of glass and ceramics via solution or sol-gel route. H. Dislich prepared a pyrex-type borosilicate glass lens by heating a compact of metal alkoxide derived powder at temperatures as low as 650  C. R. Roy prepared a millimeter-size small piece of silica glass via sol-gel route at temperatures around 1000  C. C. Mazdiyasni et al. showed that well-sintered, dense ferroelectric ceramics can be obtained at temperatures as low as 900  C, when sol-gel powders prepared from solutions of metal alkoxides are employed for sintering. Those works stimulated people’s interest in sol-gel preparation of inorganic materials, such as glasses and ceramics. Materials scientists and engineers paid attention to the possibility of this method in giving shaped materials directly from a solution without passing through the powder processing and the fact that the maximum temperature required for processing is very low compared with conventional technology for preparing glasses and ceramics. Thus, many efforts have been made in preparing bulk bodies, coating films, membranes, fibers and particles, and many commercial products were born. The significant characteristics unique to the sol-gel method became evident, when organic-inorganic hybrid materials were prepared by H. Schmidt and silica materials containing functional organic molecules were prepared by D. Avnir in the early 1980s. Such materials are produced at low temperatures near room temperature, where no decomposition of organic matter takes place. Low temperature synthesis and preparation of materials is the world of chemists. Therefore, the sol-gel method was propagated to the wide area including not only glasses and ceramics, but also organic and biomaterials. In 1990, an excellent book entitled “Sol-Gel Science” was written by Brinker and Scherer, obtaining a very high reputation. However, the remarkable scientific and technological development and broadening in the sol-gel field, together with an enormous increase in sol-gel population, appeared to demand publication of a new, comprehensive Handbook on sol-gel science and technology. ix

x

Preface to the First Edition

Thus, it was planned to publish the present Handbook, which consists of the following three volumes: • Volume 1: Sol-Gel Processing Volume editor: Prof. Hiromitsu Kozuka • Volume 2: Characterization and Properties of Sol-Gel Materials and Products Volume editor: Prof. Rui M. Almeida • Volume 3: Application of Sol-Gel Technology Volume editor: Prof. Sumio Sakka Volume 1 compiles the articles describing various aspects of sol-gel processing. Considering that the sol-gel method is a method for preparing materials, the knowledge on sol-gel processing is of primary importance to all those who are interested in sol-gel science and technology. Articles describing processing of some particular property as well as general basics for sol-gel processing are collected. Volume 2 consists of the articles dealing with characterization and properties of sol-gel materials and products. Since materials exhibit their functional properties based on their microstructure, characterization of the structure is very important. We can produce useful materials only when processing-characterization-property relationships are worked out. This indicates the importance of the articles collected in Vol. 2. The title of Vol. 3 is “Applications of Sol-Gel Technology”. The sol-gel technology is one of the methods for producing materials and so there are many other competitive methods, whenever a particular material is planned to be produced. Therefore, for the development of this excellent technology, it is important to know the sol-gel science and technology in producing new materials as well as already achieved applications. This is the purpose of Vol. 3. Sol-gel technology is a versatile technology, making it possible to produce a wide variety of materials and to provide existing materials with novel properties. I hope this three-volume Handbook will serve as an indispensable reference book for researchers, engineers, manufacturers and students working in the field of materials. Finally, I would like to express my sincere thanks to all the authors of the articles included in the Handbook for their efforts in writing excellent articles by spending their precious time. As general editor I extend my thanks to Prof. H. Kozuka and Prof. R. Almeida for their difficult work of editing each Volume. I have to confess that this Handbook would not have been realized without enthusiastic encouragement of Mr. Gregory Franklin, senior editor at Kluwer Academic Publishers. Sumio Sakka

Acknowledgments

During the course of this project, Lisa Klein and Andrei Jitianu were supported, in part, by NSF–DMR Award #1313544, Materials World Network, SusChEM: Hybrid Sol–Gel Route to Chromate-free Anticorrosive Coating. Mario Aparicio was supported in part by Ministerio de Economía y Competitividad, Spain (PCIN2013-030). This support, which enabled several opportunities to visit each other in our respective laboratories and gave us the chance to discuss the scope of this Handbook in person, was greatly appreciated. We are grateful to the members of the Springer staff who worked with us over the course of the past 3 years. We thank Mike Luby for proposing this project, Lydia Mueller for getting this underway, and we especially thank Sylvia Blago for seeing this through to completion. We also thank Clifford Nwaeburu and Santhiya Rajarathinam for technical assistance. We want to express our gratitude to the members of our International Editorial Board: Sumio Sakka, Editor Emeritus (Japan) Michel Aegerter, retired (Switzerland) Rui M. Almeida, Instituto Superior Técnico-Lisbon (Portugal) David Avnir, Hebrew University of Jerusalem (Israel) Sara A. Bilmes, Universidad de Buenos Aires (Argentina) Mary K. Carroll, Union College (USA) Dibyendu Ganguli, Indian Ceramic Society (India) Massimo Guglielmi, University of Padova (Italy) Hiromitsu Kozuka, Kansai University (Japan) Jacques Livage, Collège de France (France) Gunnar Westin, Uppsala University (Sweden)

xi

Contents

Volume 1 Part I

Processing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1

1

History of the Sol-Gel Chemistry and Technology . . . . . . . . . . . . . Sumio Sakka

3

2

The Synthesis and Solution Stability of Alkoxide Precursors Vadim G. Kessler

....

31

3

Chemistry and Applications of Polymeric Gel Precursors . . . . . . . Valery Petrykin and Masato Kakihana

81

4

Reactions of Alkoxides Toward Nanostructured or Multicomponent Oxide Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Kazumi Kato

113

5

Sol-Gel Processing of Thin Films with Metal Salts . . . . . . . . . . . . . Keishi Nishio and Toshio Tsuchiya

133

6

Aqueous Precursors Yutaka Ohya

....................................

155

7

Alkaline Silicate Solutions: An Overview of Their Structure, Reactivity, and Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Laeticia Vidal, Ameni Gharzouni, and Sylvie Rossignol

181

8

9 10

Water-Dispersed Silicates and Water-Soluble Phosphates, and Their Use in Sol-Gel Synthesis of Silicate- and Phosphate-Based Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Makoto Kobayashi, Hideki Kato, and Masato Kakihana

205

Sol-Gel Formation of Bulk Glasses . . . . . . . . . . . . . . . . . . . . . . . . . Sumio Sakka

233

....................................

257

Processing of Fibers Kanichi Kamiya

xiii

xiv

Contents

11

Stress Evolution and Cracking in Sol-Gel-Derived Thin Films . . . Hiromitsu Kozuka

275

12

Radiative Striations in Spin-Coating Films Hiromitsu Kozuka

..................

313

13

Sol-Gel Processing of Fluoride and Oxyfluoride Materials Shinobu Fujihara

......

333

14

Fluorolytic Sol-Gel Processes Erhard Kemnitz

.............................

361

15

Sol-Gel Processing of Sulfide Materials Rui M. Almeida and Jian Xu

.....................

403

16

Ultrasonic Pulverization of an Aerosol: A Versatile Tool for the Deposition of Sol-Gel Thin Films . . . . . . . . . . . . . . . . . . . . . . . . . . M. Langlet

429

17

Microparticles Preparation Using Water-in-Oil Emulsion . . . . . . . Masakazu Kawashita and Toshiki Miyazaki

18

Microwave-Assisted Sol-Gel Synthesis of Metal Oxide Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . K. M. Garadkar, A. N. Kadam, and Jinsub Park

453

483

19

Electrophoretic Sol-Gel Deposition . . . . . . . . . . . . . . . . . . . . . . . . . Atsunori Matsuda and Masahiro Tatsumisago

505

20

Electrochemical Deposition of Sol-Gel Films . . . . . . . . . . . . . . . . . Liang Liu and Daniel Mandler

531

21

Ultraviolet (UV) Irradiation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hiroaki Imai

569

22

Low-Temperature Processing of Sol-Gel Thin Films in the SiO2–TiO2 Binary System . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . M. Langlet

585

23

Borosilicates Obtained by Sol-Gel Method . . . . . . . . . . . . . . . . . . . Weidong Xiang and Jiasong Zhong

621

24

Lead-Free Ferroelectric Thin Films . . . . . . . . . . . . . . . . . . . . . . . . Barbara Malič, Alja Kupec, Katarina Vojisavljević, and Tanja Pečnik

667

25

Ferrites Obtained by Sol-Gel Method . . . . . . . . . . . . . . . . . . . . . . . Sagar E. Shirsath, Danyang Wang, Santosh S. Jadhav, M. L. Mane, and Sean Li

695

Contents

xv

....

737

....................................

765

28

Processing of High Performance Catalysts . . . . . . . . . . . . . . . . . . . Akifumi Ueno

809

29

Macroporous Morphology Control by Phase Separation . . . . . . . . Kazuki Nakanishi

835

30

Hierarchical Organization in Monolithic Sol-Gel Materials Andrea Feinle, Michael S. Elsaesser, and Nicola Hüsing

.....

867

31

Formation of Ordered Mesoporous Thin Films Through Templating . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Karen J. Edler

917

26

Oxide (TiO2) Nanotubes Obtained Through Sol-Gel Method Masahide Takahashi

27

Alumina Thin Films Marianne Nofz

Volume 2 32

Aerogel Processing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thierry Woignier, Jean Phalippou, Florence Despetis, and Sylvie Calas-Etienne

33

Aerogels from Preceramic Polymers . . . . . . . . . . . . . . . . . . . . . . . . 1013 Gian Domenico Sorarù, Emanuele Zera, and Renzo Campostrini

34

Nonhydrolytic Sol-Gel Technology . . . . . . . . . . . . . . . . . . . . . . . . . 1039 André Vioux and P. Hubert Mutin

35

Pechini Processes: An Alternate Approach of the Sol-Gel Method, Preparation, Properties, and Applications . . . . . . . . . . . . 1067 Lucangelo Dimesso

36

Modified Pechini Synthesis of Oxide Powders and Thin Films Tor Olav Løveng Sunde, Tor Grande, and Mari-Ann Einarsrud

Part II

985

. . . 1089

Characterization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1119

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1121 Rui M. Almeida and Ana C. Marques

38

Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1153 Maurizio Montagna

xvi

Contents

39

Small-Angle X-ray Scattering by Nanostructured Materials . . . . . 1185 Aldo F. Craievich

40

X-ray Absorption Spectroscopy Studies on Materials Obtained by the Sol-Gel Route . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1231 Francesco d’Acapito

41

Surface Structure of Sol-Gel-Derived Materials Using X-ray Photoelectron Spectroscopy (XPS) . . . . . . . . . . . . . . . . . . . . . . . . . 1257 Diane Holland

42

Atomic-Scale Structure of Gel Materials by Solid-State NMR . . . . 1281 Mark E. Smith and Diane Holland

43

Synthesis of Non-siliceous Glasses and Their Structural Characterization by Solid-State NMR . . . . . . . . . . . . . . . . . . . . . . 1323 Hellmut Eckert

44

Structural Characterization of Hybrid Organic–Inorganic Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1375 Plinio Innocenzi, Giovanna Brusatin, Massimo Guglielmi, and Florence Babonneau

45

Porosity Measurement Kazuki Nakanishi

46

Measurements of Gas Adsorption and Permeability of Sol-Gel Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1411 Kiyoharu Tadanaga and Tsutomu Minami

47

Specific Behavior of Sol-Gel Materials in Mercury Porosimetry: Collapse and Intrusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1425 René Pirard, Christelle Alié, and Jean-Paul Pirard

48

Viscosity and Spinnability of Gelling Solutions Sumio Sakka

49

Evolution of the Mechanical Properties During the Gel–Glass Process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1487 Thierry Woignier, Florence Despetis, P. Etienne, Adil Alaoui, L. Duffours, and Jean Phalippou

50

Mechanical and Tribological Properties of the Oxide Thin Films Obtained by Sol-Gel Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1513 Carolina J. Diliegros-Godines, Francisco Javier Flores-Ruiz, Rebeca Castanedo-Pérez, Gerardo Torres-Delgado, and Esteban Broitman

51

Characterization of the Mechanical Properties of Sol-Gel Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1527 Michel A. Aegerter

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1399

. . . . . . . . . . . . . . . 1453

Contents

xvii

52

Mechanical Properties of Organic–Inorganic Hybrids . . . . . . . . . . 1547 John D. Mackenzie and Eric P. Bescher

53

Characterization of Sol-Gel Thin-Film Waveguides . . . . . . . . . . . . 1565 Giancarlo C. Righini

54

Ellipsometry of Sol-Gel Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1595 Eric Yeatman

55

Active Sol-Gel Materials, Fluorescence Spectra, and Lifetimes . . . 1607 Anna Lukowiak, Alessandro Chiasera, Andrea Chiappini, Giancarlo C. Righini, and Maurizio Ferrari

56

Nonlinear Optical Properties of Materials Derived by Sol-Gel Technology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1651 Hiroyuki Nasu

57

Ferroelectric and Piezoelectric Properties Yuhuan Xu and John D. Mackenzie

58

Characterization of Electrical Properties . . . . . . . . . . . . . . . . . . . . 1697 Jörg Pütz, Sabine Heusing, and Michel A. Aegerter

59

Application of SVET/SIET Techniques to Study Healing Processes in Coated Metal Substrates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1727 Alexandre Bastos

60

The Use of Electrochemical Techniques for the Characterization of the Corrosion Behavior of Sol-Gel-Coated Metals . . . . . . . . . . . 1783 Francesco Andreatta and Lorenzo Fedrizzi

61

Thermal Analysis on Gels, Glasses, and Powders . . . . . . . . . . . . . 1833 Maria Zaharescu, Luminita Predoana, and Jeanina Pandele-Cusu

62

Atomistic Simulation of Sol-Gel-Derived Hybrid Materials . . . . . . 1869 Thomas S. Asche, Mirja Duderstaedt, Peter Behrens, and Andreas M. Schneider

. . . . . . . . . . . . . . . . . . . 1665

Volume 3 Part III Applications: Mechanical, Optical, Electrical, and Electrochemical . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1903

63

The Outline of Applications of the Sol-Gel Method . . . . . . . . . . . . 1905 Sumio Sakka

64

Monolithic Porous Silica for High-Speed HPLC Kazuki Nakanishi

. . . . . . . . . . . . . . 1939

xviii

Contents

65

Aerogel Sintering: From Optical Glasses to Nuclear Waste Containment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1949 Jean Phalippou, P. Dieudonné, A. Faivre, and Thierry Woignier

66

Sol-Gel Processed Membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1971 Christian Guizard, Andre Ayral, Mihai Barboiu, and Anne Julbe

67

Silica Spherical Microparticles Applied as Spacers . . . . . . . . . . . . 2019 Tatsuhiko Adachi

68

Sol-Gel Abrasive Grains: History, Precursor Properties, and Microstructural Control . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2031 Thomas E. Wood, Dwight D. Erickson, Mark G. Schwabel, and Chris J. Goodbrake

69

Sol-Gel-Prepared Glass and Ceramic Fibers . . . . . . . . . . . . . . . . . 2049 Kanichi Kamiya

70

Sol-Gel Materials for Optics and Electrooptics . . . . . . . . . . . . . . . 2065 Marcos Zayat, David Almendro, Virginia Vadillo, and David Levy

71

Sol-Gel Processed Lasers and Related Optical Materials . . . . . . . . 2093 Renata Reisfeld

72

Photonic Crystals Fabricated by Sol-Gel Process . . . . . . . . . . . . . . 2127 Makoto Kuwabara

73

Colored Coatings with Metal Colloids . . . . . . . . . . . . . . . . . . . . . . 2161 Martin Mennig and Helmut Schmidt

74

Sol-Gel Nano-/Micropatterning Process . . . . . . . . . . . . . . . . . . . . . 2177 Atsunori Matsuda and Go Kawamura

75

Sol-Gel Preparation of Reflective Coatings . . . . . . . . . . . . . . . . . . . 2205 Kensuke Makita

76

Sol-Gel-Prepared Antireflective Coatings . . . . . . . . . . . . . . . . . . . . 2223 Seiji Yamazaki

77

Sol-Gel Coatings Applied to Automotive Windows . . . . . . . . . . . . 2239 Takashige Yoneda, Sanada Yasuhiro, and Take Morimoto

78

Sol-Gel Coating of the Panel of Cathode Ray Tube for Improving the Quality of Display . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2255 Takeo Ito

79

Sol-Gel-Doped Glasses for Scintillators Masanori Koshimizu

. . . . . . . . . . . . . . . . . . . . . 2273

Contents

xix

80

Sol-Gel-Derived SnO2-Based Photonic Systems . . . . . . . . . . . . . . . 2301 Lidia Zur, Lam Thi Ngoc Tran, Marcello Meneghetti, and Maurizio Ferrari

81

Sol-Gel Processing for Spectral Hole-Burning Materials . . . . . . . . 2321 Masayuki Nogami

82

Graphene and Carbon Dots in Mesoporous Materials . . . . . . . . . . 2339 Luca Malfatti, Davide Carboni, and Plinio Innocenzi

83

Sol-Gel Protective Coatings for Metals . . . . . . . . . . . . . . . . . . . . . . 2369 Alicia Durán, Yolanda Castro, Ana Conde, and Juan José de Damborenea

84

Sol-Gel Coatings with Nanocontainers of Corrosion Inhibitors for Active Corrosion Protection of Metallic Materials . . . . . . . . . . . . . 2435 K. A. Yasakau, M. G. S. Ferreira, and Mikhail L. Zheludkevich

85

Corrosion Protection of Magnesium Alloys by Functional Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2473 Lénia M. Calado and M. F. Montemor

86

Metal Nanoparticle–Mesoporous Oxide Nanocomposite Thin Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2507 Paula C. Angelomé and M. Cecilia Fuertes

87

Xerogels, Aerogels, and Aerogel/Mineral Composites for CO2 Sequestration . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2535 Luis Esquivias, Víctor Morales-Flórez, and Alberto Santos

88

Solar Cells Based on Sol-Gel Films . . . . . . . . . . . . . . . . . . . . . . . . . 2555 Michael Grätzel

89

Sol-Gel Processing for Battery and Fuel Cell Applications Lisa Klein, Mario Aparicio, and Francoise Damay

90

Lithium Intercalation Materials for Battery Prepared by Sol-Gel Method . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2595 Jadra Mosa and Mario Aparicio

91

Sol-Gel Processing of Solid Electrolytes for Li-Ion Batteries . . . . . 2631 Nataly Carolina Rosero-Navarro and Kiyoharu Tadanaga

92

Proton Conduction in Sol-Gel-Derived Glasses and Thin Films Yusuke Daiko and Yuji Iwamoto

. . . . . . 2573

. . . 2649

xx

Contents

93

Sol-Gel-Derived Silicate-Based Composite Electrode . . . . . . . . . . . 2663 Ovadia Lev, D. Rizkov, S. Mizrahi, I. Ekeltchik, Z. G. Kipervaser, V. Gitis, A. Goifman, D. Tessema, A. Kamyshny Jr., A. D. Modestov, and J. Gun

94

Sol-Gel-Processed Photocatalytic Titania Films . . . . . . . . . . . . . . . 2695 Naoya Yoshida and Toshiya Watanabe

95

Coatings with Photocatalyst on Architectural Glass Hirokazu Tanaka and Shigeki Obana

96

Sol-Gel Coatings for Electrochromic Devices . . . . . . . . . . . . . . . . . 2745 Sabine Heusing and Michel A. Aegerter

. . . . . . . . . . . 2729

Volume 4 Part IV Applications: Preservation, Organic–Inorganic Hybrids, and Bio-related Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

2793

97

Sol-Gel Wood Preservation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2795 Thomas Hübert and Muhammad Shabir Mahr

98

Sol-Gel Materials for Art Conservation . . . . . . . . . . . . . . . . . . . . . 2843 Eric P. Bescher and John D. Mackenzie

99

Sol-Gel Science and Cultural Heritage . . . . . . . . . . . . . . . . . . . . . . 2859 George Wheeler

100

Sol-Gel Environmental Sensors for Preventive Conservation of Cultural Heritage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2877 Maria-Angeles Villegas, Javier Peña-Poza, and Manuel Garcia-Heras

101

Encapsulation of Enzymes, Antibodies, and Bacteria Jacques Livage and Thibaud Coradin

102

Entrapment of Organic Molecules . . . . . . . . . . . . . . . . . . . . . . . . 2933 Kazunori Matsui, Hiromasa Nishikiori, and Tsuneo Fujii

103

Hybrid Materials for Molecular Sieves . . . . . . . . . . . . . . . . . . . . . 2973 Johan E. ten Elshof

104

Click Functionalization of Sol-Gel Materials . . . . . . . . . . . . . . . . 3001 Shridevi Shenoi-Perdoor, Achraf Noureddine, Fabien Dubois, Michel Wong Chi Man, and Xavier Cattoën

105

Sol-Gel Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3041 Massimo Guglielmi and Alessandro Martucci

. . . . . . . . . 2909

Contents

xxi

106

Hybrid Materials for Micro- and Nanofabrication . . . . . . . . . . . . 3065 Laura Brigo, Gioia Della Giustina, and Giovanna Brusatin

107

Architecture of Silsesquioxanes . . . . . . . . . . . . . . . . . . . . . . . . . . . 3119 Sandra Dirè, Evgeny Borovin, and François Ribot

108

Polyhedral Silsesquioxanes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3153 Abby R. Jennings, Scott T. Iacono, and Joseph M. Mabry

109

Mesoporous Polysilsesquioxanes: Preparation, Properties, and Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3177 Douglas A. Loy

110

Hybrid Nanocomposites Through Colloidal Interactions Between Crystalline Polysaccharide Nanoparticles and Oxide Precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3213 Emmanuel Belamie and Bruno Alonso

111

Anti-soiling Effect of Porous SiO2 Coatings . . . . . . . . . . . . . . . . . 3253 Peer Löbmann

112

Sol-Gel Preparation of Crystalline Oxide Thin Films on Plastics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3271 Hiromitsu Kozuka

113

Functional Barrier Coatings on the Basis of Hybrid Polymers . . . 3295 Sabine Amberg-Schwab

114

Hybrid Aerogels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3317 Kazuyoshi Kanamori

115

Carbon Aerogels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3339 Marcus A. Worsley and Theodore F. Baumann

116

Mechanical Behavior of Nanocomposite Aerogels Thierry Woignier, Juan Primera, Adil Alaoui, and Sylvie Calas-Etienne

117

The Development of Quantum Dot/Silica Particles for Fluorescence Imaging and Medical Diagnostics . . . . . . . . . . . . . . 3393 Yoshio Kobayashi and Kohsuke Gonda

118

Hybrid Sol-Gels for DNA Arrays and Other Lab-on-a-Chip Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3431 Caner Durucan and Carlo G. Pantano

119

Antimicrobial Coatings Obtained by Sol-Gel Method . . . . . . . . . 3461 Boris Mahltig, Thomas Grethe, and Hajo Haase

. . . . . . . . . . . . 3375

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Contents

120

Bioactive Ceramic Porcelain/Glass for Dental Application . . . . . 3489 E. Kontonasaki, X. Chatzistavrou, K. M. Paraskevopoulos, and P. Koidis

121

Bioactive Silica-Based Coating on Stainless Steel Implants . . . . . 3505 Josefina Ballarre and Silvia M. Ceré

122

Biomaterials Obtained by Gelation . . . . . . . . . . . . . . . . . . . . . . . . 3555 Alain C. Pierre

123

Sol-Gel Silica-Based Biomaterials and Bone Tissue Regeneration . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3597 María Vallet-Regí and Antonio J. Salinas

124

Inorganic-Organic Hybrids for Biomedical Applications . . . . . . . 3619 Yuki Shirosaki, Yuri Nakamura, Tomohiko Yoshioka, and Akiyoshi Osaka

125

Enzymatic Sol-Gel Biosensors . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3705 Elena Casero, M. D. Petit-Domínguez, and Luis Vázquez

Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3745

About the Editors

Dr. Lisa Klein obtained a BS in Metallurgy in 1973 and a Ph.D. in Ceramics in 1977 from the Material Science and Engineering Department at the Massachusetts Institute of Technology. With Rutgers since 1977, she became a full Professor in 1987. In 1998, she received the Achievement Award of the Society of Women Engineers for “breakthrough contributions in sol-gel science and engineering, particularly sol-gel applications in electrolytes, electrochromics, membranes and nanocomposites.” Professor Klein’s areas of expertise are sol-gel processing of glasses, ceramics, and organicinorganic hybrids. Her use of the sol-gel process finds applications in planar waveguides, ceramic membranes, solid electrolytes, and components for fuel cells. In 2017, she received the “Life Achievement Award” from the International Sol-Gel Society, in Liege, Belgium. Dr. Mario Aparicio completed his Ph.D. in 1998 at the Universidad Autónoma de Madrid, Spain. He studied chemistry, mainly focusing on sol-gel processing and ceramic matrix composites. In 1999, he was awarded a 2-year fellowship from Ministerio de Educación y Cultura to study hybrid proton exchange membranes at Rutgers University, Piscataway, NJ. Presently, he is a researcher in the Instituto de Cerámica y Vidrio, CSIC, Madrid, where he continues to develop PEM membranes, electrodes and electrolytes for batteries, and sol-gel corrosion protection coatings.

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About the Editors

Dr. Andrei Jitianu completed his Ph.D. in 2001 at the University of Bucharest, Romania. He studied materials chemistry, mainly focusing on sol-gel processes. Since then, he has done postdoctoral research at the University of Orléans, France; Clarkson University, Potsdam, NY; and Rutgers University, Piscataway, NJ. In 2008, he joined The City University of New York, Lehman College, Bronx, NY, where he became a full Professor in 2017. He has experience with sol-gel inorganic-organic hybrids for a variety of applications. In 2005, he received “Gh. Spacu” Award of the Romanian Academy for “Chemistry of the Sol-Gel Processes in Oxide and Hybrid Systems.” In 2010, he received the Feliks Gross Endowment Award, as one of only two junior faculties selected for their research excellence from the Graduate Center of the entire City University of New York. Professor Jitianu’s areas of expertise are sol-gel processing of oxide and hybrid systems, coatings, “melting gels,” protection against corrosion, nanomaterials for catalysis, and layer double hydroxides.

Contributors

Tatsuhiko Adachi UBE EXSYMO Co., Ltd., Chuo-ku, Tokyo, Japan Michel A. Aegerter Bottens, Switzerland INM – Leibniz Institute for New Materials, Saarbruecken, Germany Adil Alaoui Faculté des Sciences et Techniques de Tanger, Tanger, Morocco Christelle Alié Department of Chemical Engineering, School of Engineering, University of Liège, Liège, Belgium Rui M. Almeida Departamento de Engenharia Química / Centro de Química Estrutural (CQE), Instituto Superior Técnico / Universidade de Lisboa, Lisboa, Portugal David Almendro Sol-Gel Group – SGG, Instituto de Ciencia de Materiales de Madrid, ICMM. CSIC, Madrid, Spain Bruno Alonso Institut Charles Gerhardt Montpellier, UMR 5253 CNRS/UM/ ENSCM, ENSCM, Montpellier, France Ecole Pratique des Hautes Etudes, PSL Research University, Paris, France Sabine Amberg-Schwab Fraunhofer-Institut für Silicatforschung ISC, Würzburg, Germany Francesco Andreatta Polytechnic Department of Engineering and Architecture, University of Udine, Udine, Italy Paula C. Angelomé Gerencia Química, Centro Atómico Constituyentes, Comisión Nacional de Energía Atómica, CONICET, San Martín, Buenos Aires, Argentina Mario Aparicio Instituto de Cerámica y Vidrio, Consejo Superior de Investigaciones Científicas (CSIC), Madrid, Spain Thomas S. Asche Institut für Anorganische Chemie, Leibniz Universität Hannover, Hannover, Germany Andre Ayral Universite Montpellier 2, Institut Européen des Membranes, Montpellier, Cedex 5, France Florence Babonneau l’Institut des Matériaux de Paris-Centre, Paris, France xxv

xxvi

Contributors

Josefina Ballarre Materials Science and Technology Research Institute (INTEMA), University of Mar del Plata – National Research and Technology Council (CONICET), Mar del Plata, Argentina Mihai Barboiu Institut European des Membranes, IEM – CNRS 5635, Montpellier, France Alexandre Bastos DEMaC – Department of Materials and Ceramic Engineering, CICECO – Aveiro Institute of Materials, University of Aveiro, Aveiro, Portugal Theodore F. Baumann Physical and Life Sciences Directorate, Lawrence Livermore National Laboratory, Livermore, CA, USA Peter Behrens Institut für Anorganische Chemie, Leibniz Universität Hannover, Hannover, Germany Emmanuel Belamie Institut Charles Gerhardt Montpellier, UMR 5253 CNRS/ UM/ENSCM, ENSCM, Montpellier, France Eric P. Bescher Department of Materials Science and Engineering, University of California, Los Angeles, Los Angeles, CA, USA Evgeny Borovin Dipartimento di Ingegneria Industriale, Università di Trento, Trento, Italy Laura Brigo Department of Industrial Engineering, University of Padova, Padova, Italy Center for Materials and Microsystems, Bruno Kessler Foundation, Trento, Italy Esteban Broitman Thin Film Physics Division, Linköping University, Linköping, Sweden Giovanna Brusatin Department of Industrial Engineering, University of Padova, Padova, Italy Lénia M. Calado CQE, DEQ, Instituto Superior Técnico, Universidade de Lisboa, Lisbon, Portugal Sylvie Calas-Etienne Laboratoire Charles Coulomb, Université Montpellier 2, Montpellier Cedex 5, France Renzo Campostrini Department of Industrial Engineering, University of Trento, Trento, Italy Davide Carboni Laboratory of Materials Science and Nanotechnology, LMNT – D.A.D.U., University of Sassari and CR-INSTM, Alghero, Sassari, Italy Elena Casero Departamento de Química Analítica y Análisis Instrumental, Facultad de Ciencias, Campus de Excelencia de la Universidad Autónoma de Madrid, Madrid, Spain

Contributors

xxvii

Rebeca Castanedo-Pérez CINVESTAV-Unidad Querétaro, Querétaro, Qro, Mexico Yolanda Castro Instituto de Cerámica y Vidrio (CSIC), Madrid, Spain Xavier Cattoën University of Grenoble Alpes, Inst NEEL, Grenoble, France CNRS, Institut NEEL, Grenoble, France Silvia M. Ceré Materials Science and Technology Research Institute (INTEMA), University of Mar del Plata – National Research and Technology Council (CONICET), Mar del Plata, Argentina X. Chatzistavrou Department of Orthodontics and Pediatric Dentistry, School of Dentistry, University of Michigan, Ann Arbor, MI, USA Andrea Chiappini IFN-CNR CSMFO Lab, and FBK Photonics Unit, Trento, Italy Alessandro Chiasera IFN-CNR CSMFO Lab, and FBK Photonics Unit, Trento, Italy Ana Conde Centro Nacional de Investigaciones Metalúrgicas (CSIC), Madrid, Spain Thibaud Coradin Chaire de Chimie de la Matière Condensée, Collège de France, Paris, France Aldo F. Craievich Institute of Physics, University of São Paulo, São Paulo, Brasil Francesco d’Acapito CNR-IOM-OGG c/o European Synchrotron Radiation Facility, LISA CRG, Grenoble, France Yusuke Daiko Department of Frontier Materials, Nagoya Institute of Technology, Nagoya, Aichi, Japan Francoise Damay Laboratoire Leon Brillouin, CEA-Saclay Bat. 563, Gif-surYvette, France Juan José de Damborenea Centro Nacional de Investigaciones Metalúrgicas (CSIC), Madrid, Spain Gioia Della Giustina Department of Industrial Engineering, University of Padova, Padova, Italy Florence Despetis Laboratoire Charles Coulomb, Université Montpellier 2, Montpellier Cedex 5, France P. Dieudonné Laboratoire Charles Coulomb, Université Montpellier 2, Montpellier Cedex 5, France Carolina J. Diliegros-Godines CINVESTAV-Unidad Querétaro, Querétaro, Qro, Mexico CNyN-UNAM, Ensenada, BC, Mexico

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Contributors

Lucangelo Dimesso Earth and Material Sciences Department, Technische Universitaet Darmstadt, Darmstadt, Germany Sandra Dirè Dipartimento di Ingegneria Industriale, Università di Trento, Trento, Italy Fabien Dubois University of Grenoble Alpes, Inst NEEL, Grenoble, France CNRS, Institut NEEL, Grenoble, France Mirja Duderstaedt Institut für Anorganische Chemie, Leibniz Universität Hannover, Hannover, Germany L. Duffours PRIME Verre, Montpellier, France Le Lamentin, Martinique, France Alicia Durán Instituto de Cerámica y Vidrio, Madrid, Spain Caner Durucan Department of Metallurgical and Materials Engineering, Middle East Technical University, Ankara, Turkey Hellmut Eckert Institut für Physikalische Chemie, Westfälische Wilhelms Universität Münster, Münster, Germany Instituto de Física, São Carlos, Universidade de São Paulo, São Carlos, SP, Brazil Karen J. Edler Department of Chemistry, University of Bath, Bath, UK Mari-Ann Einarsrud Department of Materials Science and Engineering, NTNU Norwegian University of Science and Technology, Trondheim, Norway I. Ekeltchik The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel Michael S. Elsaesser Department Chemistry and Physics of Materials, Paris Lodron University Salzburg, Salzburg, Austria Johan E. ten Elshof MESA+ Institute for Nanotechnology, University of Twente, Enschede, The Netherlands Dwight D. Erickson Minnesota Mining and Manufacturing Company, Stillwater, MN, USA Luis Esquivias Departamento Física de la Materia Condensada, Universidad de Sevilla, Instituto de Ciencia de Materiales de Sevilla (CSIC-US), Sevilla, Spain P. Etienne Laboratoire Charles Coulomb, Montpellier Université, Montpellier, France A. Faivre Laboratoire Charles Coulomb, Montpellier, France Lorenzo Fedrizzi Polytechnic Department of Engineering and Architecture, University of Udine, Udine, Italy

Contributors

xxix

Andrea Feinle Department Chemistry and Physics of Materials, Paris Lodron University Salzburg, Salzburg, Austria Maurizio Ferrari Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi, Rome, Italy CNR-IFN, CSMFO Lab. and FBK Photonics Unit, Trento, Italy M. G. S. Ferreira Department of Materials and Ceramic Engineering, CICECO – Aveiro Institute of Materials, University of Aveiro, Aveiro, Portugal Francisco Javier Flores-Ruiz CINVESTAV-Unidad Querétaro, Querétaro, Qro, Mexico CNyN-UNAM, Ensenada, BC, Mexico M. Cecilia Fuertes Gerencia Química, Centro Atómico Constituyentes, Comisión Nacional de Energía Atómica, CONICET, San Martín, Buenos Aires, Argentina Shinobu Fujihara Department of Applied Chemistry, Faculty of Science and Technology, Keio University, Yokohama, Japan Tsuneo Fujii Nagano Prefectural Institute of Technology, Ueda, Japan K. M. Garadkar Nanomaterials Research Laboratory, Department of Chemistry, Shivaji University, Kolhapur, Maharashtra, India Manuel Garcia-Heras Consejo Superior de Investigaciones Cientificas (CSIC), Instituto de Historia, CCHS, Madrid, Spain Ameni Gharzouni SPCTS, University of Limoges, Limoges, France V. Gitis The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel A. Goifman The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel Kohsuke Gonda Department of Health Sciences, Graduate School of Medicine, Tohoku University, Sendai, Miyagi, Japan Chris J. Goodbrake Minnesota Mining and Manufacturing Company, Stillwater, MN, USA Tor Grande Department of Materials Science and Engineering, NTNU Norwegian University of Science and Technology, Trondheim, Norway Michael Grätzel Laboratory of Photonics and Interfaces, Institute of Chemical Science and Engineering, Ecole Polytechnique Fédérale de Lausanne (EPFL), Lausanne, Switzerland

xxx

Contributors

Thomas Grethe Research Institute for Textile and Clothing, Niederrhein University of Applied Sciences, Mönchengladbach, Germany Massimo Guglielmi Dipartimento di Ingegneria Industriale, Università di Padova, Padova, Italy Christian Guizard L’Institut Européen des Membranes, Montpellier, France J. Gun The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel Hajo Haase Institut für Lebensmitteltechnologie und Lebensmittelchemie, Technische Universität Berlin, Berlin, Germany Sabine Heusing Campus D2 2, INM – Leibniz Institute for New Materials, Saarbruecken, Germany Diane Holland Department of Physics, University of Warwick, Coventry, UK P. Hubert Mutin Institut Charles Gerhardt, UMR5253 CNRS-UM-ENSCM, Université de Montpellier, Montpellier, France Thomas Hübert Federal Institute for Materials Research and Testing (BAM), Berlin, Germany Nicola Hüsing Department Chemistry and Physics of Materials, Paris Lodron University Salzburg, Salzburg, Austria Scott T. Iacono Department of Chemistry and Chemistry Research Center, United States Air Force Academy, Colorado Springs, CO, USA Hiroaki Imai Department of Applied Chemistry, Faculty of Science and Technology, Keio University, Yokohama, Turkey Plinio Innocenzi Laboratory of Materials Science and Nanotechnology, LMNT – D.A.D.U., University of Sassari and CR-INSTM, Alghero, Sassari, Italy Takeo Ito Toshiba Corporation, Kumagaya-shi, Saitama, Japan Yuji Iwamoto Department of Frontier Materials, Nagoya Institute of Technology, Nagoya, Aichi, Japan Santosh S. Jadhav Department of Physics, Dnyanopasak Shikshan Mandal’s Arts, Commerce and Science College, Jintur, India Abby R. Jennings Department of Chemistry and Chemistry Research Center, United States Air Force Academy, Colorado Springs, CO, USA Anne Julbe Universite Montpellier 2, Institut Européen des Membranes, Montpellier, Cedex 5, France A. N. Kadam Department of Electronics and Computer Engineering, Hanyang University, Seoul, Republic of Korea

Contributors

xxxi

Masato Kakihana Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, Sendai, Miyagi, Japan Kanichi Kamiya Department of Industrial Chemistry, Mie University, Tsu, Japan A. Kamyshny Jr. The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel Kazuyoshi Kanamori Department of Chemistry, Graduate School of Science, Kyoto University, Kyoto, Japan Kazumi Kato National Institute of Advanced Industrial Science and Technology (AIST), Inorganic Functional Materials Research Institute (IFMRI), Nagoya, Japan Hideki Kato Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, Sendai, Miyagi, Japan Go Kawamura Department of Electrical and Electronic Information Engineering, Toyohashi University of Technology, Toyohashi, Aichi, Japan Masakazu Kawashita Graduate School of Biomedical Engineering, Tohoku University, Sendai, Miyagi, Japan Erhard Kemnitz Chemistry Department, Humboldt-Universität zu Berlin, Berlin, Germany Vadim G. Kessler Department of Chemistry and Biotechnology, Swedish University of Agricultural Sciences (SLU), Uppsala, Sweden Z. G. Kipervaser The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel Lisa Klein Materials Science and Engineering, Rutgers University, Piscataway, NJ, USA Makoto Kobayashi Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, Sendai, Miyagi, Japan Yoshio Kobayashi Department of Biomolecular Functional Engineering, College of Engineering, Ibaraki University, Hitachi, Ibaraki, Japan P. Koidis Department of Fixed Prosthesis and Implant Prosthodontics, Faculty of Dentistry, School of Health Sciences, Aristotle University of Thessaloniki, Thessaloniki, Greece E. Kontonasaki Department of Fixed Prosthesis and Implant Prosthodontics, Faculty of Dentistry, School of Health Sciences, Aristotle University of Thessaloniki, Thessaloniki, Greece Masanori Koshimizu Department of Applied Chemistry, Graduate School of Engineering, Tohoku University, Sendai, Miyagi, Japan

xxxii

Contributors

Hiromitsu Kozuka Department of Chemistry and Materials Engineering, Faculty of Chemistry, Materials and Bioengineering, Kansai University, Osaka, Japan Alja Kupec Electronic Ceramics Department, Jožef Stefan Institute, Ljubljana, Slovenia Faculty of Mechanical Engineering, Laboratory for Tribology and Interface Nanotechnology, University of Ljubljana, Ljubljana, Slovenia Makoto Kuwabara Kyushu University, Kasuga, Fukuoka, Japan M. Langlet Grenoble INP – Minatec, Grenoble, France Ovadia Lev The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel The Casali Center of Applied Chemistry, The Chemistry Institute, The Hebrew University of Jerusalem, Jerusalem, Israel David Levy Sol-Gel Group – SGG, Instituto de Ciencia de Materiales de Madrid, ICMM. CSIC, Madrid, Spain Sean Li School of Materials Science and Engineering, University of New South Wales, Sydney, NSW, Australia Liang Liu Laboratoire de Chimie Physique et Microbiologie pour l’Environnement, UMR 7564, CNRS-Université de Lorraine, Villers-lès-Nancy, France Jacques Livage Chaire de Chimie de la Matière Condensée, Collège de France, Paris, France Peer Löbmann Optik und Elektronik, Fraunhofer-Institut für Silicatforschung ISC, Würzburg, Germany Douglas A. Loy Department of Materials Science and Engineering, Department of Chemistry and Biochemistry, The University of Arizona, Tucson, AZ, USA Anna Lukowiak Institute of Low Temperature and Structure Research PAS, Wroclaw, Poland Joseph M. Mabry Aerospace Systems Directorate, Air Force Research Laboratory, Edwards AFB, CA, USA John D. Mackenzie Department of Materials Science and Engineering, University of California, Los Angeles, Los Angeles, CA, USA Boris Mahltig Niederrhein University of Applied Sciences, Faculty of Textile and Clothing Technology, Mönchengladbach, Germany Kensuke Makita Department of Materials Science and Engineering, Nagoya Institute of Technology, Nagoya, Japan

Contributors

xxxiii

Luca Malfatti Laboratory of Materials Science and Nanotechnology, LMNT – D.A.D.U., University of Sassari and CR-INSTM, Alghero, Sassari, Italy Barbara Malič Electronic Ceramics Department, Jožef Stefan Institute, Ljubljana, Slovenia Daniel Mandler Institute of Chemistry, The Hebrew University of Jerusalem, Jerusalem, Israel M. L. Mane Department of Physics, SGRG Shinde Mahavidyalaya, Osmanabad, Maharashtra, India Ana C. Marques Departamento de Engenharia Química / Centre for Natural Resources and the Environment (CERENA), Instituto Superior Técnico, Universidade de Lisboa, Lisboa, Portugal Alessandro Martucci Dipartimento di Ingegneria Industriale, Università di Padova, Padova, Italy Atsunori Matsuda Department of Electrical and Electronic Information Engineering, Toyohashi University of Technology, Toyohashi, Aichi, Japan Kazunori Matsui Graduate School of Engineering, Kanto Gakuin University, Yokohama, Kanagawa, Japan Marcello Meneghetti CNR-IFN, CSMFO Lab. and FBK Photonics Unit, Trento, Italy Department of Physics, University of Trento, Trento, Italy Martin Mennig Department of Glasses, Ceramics and Composites, Glass Group, Institut für Neue Materialien, Saarbrücken, Germany Tsutomu Minami Osaka Prefecture University, Osaka, Japan Toshiki Miyazaki Graduate School of Life Science and Systems Engineering, Kyushu Institute of Technology, Kitakyushu-shi, Fukuoka, Japan S. Mizrahi The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel A. D. Modestov The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel Maurizio Montagna Disordered System Laboratory, Department of Physics, Università degli Studi di Trento, Trento, Italy M. F. Montemor CQE, DEQ, Instituto Superior Técnico, Universidade de Lisboa, Lisbon, Portugal

xxxiv

Contributors

Víctor Morales-Flórez Departamento Física de la Materia Condensada, Universidad de Sevilla, Sevilla, Spain Take Morimoto Morimoto Co., Ltd., Fujimi-shi, Saitama, Japan Jadra Mosa Consejo Superior de Investigaciones Científicas (CSIC), Instituto de Cerámica y Vidrio (ICV), Madrid, Spain Yuri Nakamura Department of Applied Chemistry and Biotechnology, Okayama University, Okayama-shi, Japan Kazuki Nakanishi Department of Chemistry, Graduate School of Chemistry, Kyoto University, Kyoto, Japan Hiroyuki Nasu Division of Chemistry for Materials, Graduate School of Engineering, Mie University, Tsu, Japan Hiromasa Nishikiori Department of Materials Chemistry, Faculty of Engineering, Shinshu University, Nagano, Japan Keishi Nishio Tokyo University of Science, Materials Science and Technology, Noda-shi, Japan Marianne Nofz Division 5.6 “Glas”, Bundesanstalt für Materialforschung und – prüfung, Berlin, Germany Masayuki Nogami Toyota Physical and Chemical Research Institute, Nagakute, Aichi, Japan Achraf Noureddine Institut Charles Gerhardt Montpellier, UMR-5253, CNRS, ENSCM, Université Montpellier, Montpellier, France Shigeki Obana Nippon Sheet Glass Co. Ltd., Tokyo, Japan Yutaka Ohya Department of Chemistry and Biomolecular Science, Gifu University, Gifu, Japan Akiyoshi Osaka Faculty of Engineering, Okayama University, Okayama-shi, Japan Jeanina Pandele-Cusu “Ilie Murgulescu” Institute of Physical Chemistry of the Romanian Academy, Bucharest, Romania Carlo G. Pantano Department of Materials Science and Engineering, Materials Research Institute, The Pennsylvania State University, University Park, PA, USA K. M. Paraskevopoulos Department of Physics, Faculty of Natural Sciences, Aristotle University of Thessaloniki, Thessaloniki, Greece Jinsub Park Department of Electronics and Computer Engineering, Hanyang University, Seoul, Republic of Korea Tanja Pečnik Electronic Ceramics Department, Jožef Stefan Institute, Ljubljana, Slovenia

Contributors

xxxv

Javier Peña-Poza Consejo Superior de Investigaciones Cientificas (CSIC), Instituto de Historia, CCHS, Madrid, Spain M. D. Petit-Domínguez Departamento de Química Analítica y Análisis Instrumental, Facultad de Ciencias, Campus de Excelencia de la Universidad Autónoma de Madrid, Madrid, Spain Valery Petrykin Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, Sendai, Miyagi, Japan Jean Phalippou Laboratoire Charles Coulomb, Université Montpellier 2, Montpellier Cedex 5, France Alain C. Pierre Institut de recherches sur la Catalyse et l’Environnement de Lyon, CNRS, UMR 5256, Université Claude Bernard Lyon 1, Université Claude Bernard Lyon 1, Villeurbanne, France Jean-Paul Pirard Department of Applied Chemistry, School of Engineering, University of Liège, Liège, Belgium René Pirard Laboratoire de Génie Chimique, Institutde Chimie B6a, University of Liège, Liège, Belgium Luminita Predoana “Ilie Murgulescu” Institute of Physical Chemistry of the Romanian Academy, Bucharest, Romania Juan Primera Departamento de Fisica, FEC, LUZ, Maracaibo, Venezuela Escuela Superior Politécnica del Litoral (ESPOL) Facultad de Ciencias Naturales y Matemáticas, Departamento de Física, Campus Gustavo Galindo, Guayaquil, Ecuador Jörg Pütz Carl Zeiss Smart Optics GmbH, Aalen, Germany Renata Reisfeld The Hebrew University of Jerusalem, Jerusalem, Israel François Ribot Laboratoire de Chimie de la Matière Condensée de Paris (LCMCP), Sorbonne Universités, UPMC Université Paris 06, CNRS, Collège de France, Paris, France Giancarlo C. Righini Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi, Rome, Italy Nello Carrara Institute of Applied Physics, CNR, Sesto Fiorentino – Firenze, Italy D. Rizkov The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel

Alain C. Pierre has retired.

xxxvi

Contributors

Nataly Carolina Rosero-Navarro Faculty of Engineering, Hokkaido University, Hokkaido, Japan Sylvie Rossignol Ecole Nationale Superieure de Ceramique Industrielle–Limoges, Limoges, France Sumio Sakka Sakka Laboratory, Kuzuha-Asahi, Hirakata, Osaka-fu, Japan Antonio J. Salinas Departamento de Quimica Inorgánica y Bioinorgánica, Universidad Complutense de Madrid, Instituto de Investigación Sanitaria Hospital, Madrid, Spain Networking Research Center on Bioengineering, Biomaterials and Nanomedicine (CIBER-BBN), Madrid, Spain Alberto Santos Departamento Ciencias de la Tierra, Universidad de Cádiz, Puerto Real, Cádiz, Spain Helmut Schmidt HSM TechConsult GmbH, Saarbrücken, Germany Andreas M. Schneider Institut für Anorganische Chemie, Leibniz Universität Hannover, Hannover, Germany Mark G. Schwabel Minnesota Mining and Manufacturing Company, Stillwater, MN, USA Muhammad Shabir Mahr Federal Institute for Materials Research and Testing (BAM), Berlin, Germany Shridevi Shenoi-Perdoor University of Grenoble Alpes, Inst NEEL, Grenoble, France CNRS, Institut NEEL, Grenoble, France Yuki Shirosaki Graduate School of Engineering, Kyushu Institute of Technology, Kitakyushu, Japan Sagar E. Shirsath School of Materials Science and Engineering, University of New South Wales, Sydney, NSW, Australia Mark E. Smith Lancaster University, Lancaster, UK Gian Domenico Sorarù Department of Industrial Engineering, University of Trento, Trento, Italy Tor Olav Løveng Sunde Department of Sustainable Energy Technology, SINTEF Materials and Chemistry, Oslo, Norway Kiyoharu Tadanaga Faculty of Engineering, Hokkaido University, Hokkaido, Japan Masahide Takahashi Division of Materials Science, Graduate School of Engineering, Osaka Prefecture University, Osaka, Japan Hirokazu Tanaka Nippon Sheet Glass Co. Ltd., Tokyo, Japan

Contributors

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Masahiro Tatsumisago Department of Applied Chemistry, Osaka Prefecture University, Sakai, Osaka, Japan D. Tessema The Casali Center of Applied Chemistry and the Harvey M. Kruger Center for Nanoscience and Nanotechnology, The Hebrew University of Jerusalem, Jerusalem, Israel Gerardo Torres-Delgado CINVESTAV-Unidad Querétaro, Querétaro, Qro, Mexico Lam Thi Ngoc Tran CNR-IFN, CSMFO Lab. and FBK Photonics Unit, Trento, Italy Department of Civil, Environmental and Mechanical Engineering, University of Trento, Trento, Italy Ho Chi Minh City University of Technical Education, Ho Chi Minh City, Viet Nam Toshio Tsuchiya Nagasaki University, Nagasaki, Japan Akifumi Ueno Department of Materials Science, Faculty of Engineering, Shizuoka University, Hamamatsu, Japan Virginia Vadillo Sol-Gel Group – SGG, Instituto de Ciencia de Materiales de Madrid, ICMM. CSIC, Madrid, Spain María Vallet-Regí Departamento de Quimica Inorgánica y Bioinorgánica, Universidad Complutense de Madrid, Instituto de Investigación Sanitaria Hospital, Madrid, Spain Networking Research Center on Bioengineering, Biomaterials and Nanomedicine (CIBER-BBN), Madrid, Spain Luis Vázquez Instituto de Ciencia de Materiales de Madrid (CSIC), Campus de Excelencia de la Universidad Autónoma de Madrid, Madrid, Spain Laeticia Vidal SPCTS, University of Limoges, Limoges, France Maria-Angeles Villegas Consejo Superior de Investigaciones Cientificas (CSIC), Instituto de Historia, CCHS, Madrid, Spain André Vioux Institut Charles Gerhardt, UMR5253 Université de Montpellier, Montpellier, France

CNRS-UM-ENSCM,

Katarina Vojisavljević Electronic Ceramics Department, Jožef Stefan Institute, Ljubljana, Slovenia Danyang Wang School of Materials Science and Engineering, University of New South Wales, Sydney, NSW, Australia Toshiya Watanabe Policy Alternatives Research Institute, The University of Tokyo, Tokyo, Japan George Wheeler Columbia University, New York, NY, USA

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Contributors

Thierry Woignier IMBE, CNRS, IRD, Aix Marseille Université, Avignon Université, Marseille, France IRD - Campus Agro Environnemental Caraïbes, Le Lamentin, Martinique, France Michel Wong Chi Man Institut Charles Gerhardt Montpellier, UMR-5253, CNRS, ENSCM, Université Montpellier, Montpellier, France Thomas E. Wood Minnesota Mining and Manufacturing Company, Stillwater, MN, USA Marcus A. Worsley Physical and Life Sciences Directorate, Lawrence Livermore National Laboratory, Livermore, CA, USA Weidong Xiang College of Chemistry and Materials Engineering, Wenzhou University, Wenzhou, China Jian Xu Faculty of Electrical Engineering and Computer Science, Ningbo University, Ningbo, Zhejiang, China Yuhuan Xu Department of Materials Science and Engineering, University of California, Los Angeles, Los Angeles, CA, USA Seiji Yamazaki Glass Research Center, Central Glass Co. Ltd., Matsusaka, Japan K. A. Yasakau Department of Materials and Ceramic Engineering, CICECO – Aveiro Institute of Materials, University of Aveiro, Aveiro, Portugal Sanada Yasuhiro Asahi Glass Co., Ltd., Tokyo, Japan Eric Yeatman Faculty of Engineering, Department of Electrical and Electronic Engineering, Imperial College London, London, UK Takashige Yoneda Asahi Glass Co., Ltd., Tokyo, Japan Naoya Yoshida Department of Applied Chemistry, School of Advanced Engineering, Kogakuin University, Hachioji City, Tokyo, Japan Tomohiko Yoshioka Graduate School of Natural Science and Technology, Okayama University, Okayama-shi, Japan Maria Zaharescu “Ilie Murgulescu” Institute of Physical Chemistry of the Romanian Academy, Bucharest, Romania Marcos Zayat Sol-Gel Group – SGG, Instituto de Ciencia de Materiales de Madrid, ICMM. CSIC, Madrid, Spain Emanuele Zera Department of Industrial Engineering, University of Trento, Trento, Italy Mikhail L. Zheludkevich Magnesium Innovation Centre, MagIC at HelmholtzZentrum Geesthacht, Geesthacht, Germany

Contributors

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Institute for Materials Science, Faculty of Engineering, Kiel University, Kiel, Germany Jiasong Zhong College of Materials and Environmental Engineering, Hangzhou Dianzi University, Hangzhou, China Lidia Zur Museo Storico della Fisica e Centro Studi e Ricerche Enrico Fermi, Rome, Italy CNR-IFN, CSMFO Lab. and FBK Photonics Unit, Trento, Italy

Part I Processing

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History of the Sol‐Gel Chemistry and Technology Sumio Sakka

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . History of Overall Sol‐Gel Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Sol‐Gel Technology Before the 1960s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Sol‐Gel Technology in the Late 1960s to Early 1980s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . The First Sol‐Gel Workshop in 1981 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Novel Materials Concept Derived from Sol‐Gel Technology Around 1984 . . . . . . . . . . . . . . . . Development in 1985–1995: New Functional Materials by the Sol‐Gel Method . . . . . . . . . . . Progress of Sol‐Gel Technology in 1995–2005 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Progress of Sol‐Gel Technology in 2005– 2015 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . History of Dye-Sensitized Solar Cells . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Grätzel’s Dye-Sensitized Solar Cells . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Dye-Sensitized Solar Cell with Solid Electrolyte . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Perovskite Solar Cell . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . History of Sol‐Gel-Derived Ferroelectric Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Sol‐Gel Preparation of Ferroelectric Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Sol‐Gel Preparation of Ferroelectric Materials in 1980–1995 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Progress in 1995–2005 and Tasks Imposed by Restriction on Lead-Containing Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Progress in 2005–2015: Sol‐Gel Processing of Multiferroic Materials . . . . . . . . . . . . . . . . . . . . . . Recent State of the Art of Ferroelectric Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . History of Sol‐Gel Biochemical and Biomedical Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hench’s Suggestion on Compatibility of Silicates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Starting and Rapid Development of Sol‐Gel Preparation of Biochemical and Biomedical Materials in the 1990s . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Progress in 2000–2006 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Progress in 2006–2015 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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S. Sakka (*) Sakka Laboratory, Kuzuha-Asahi, Hirakata, Osaka-fu, Japan e-mail: [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_87

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History of Processing Porous Materials with Controlled Pores . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Aerogels (1931, 1968) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mesoporous Materials Based on Self-Assembly (1990, 1992) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hierarchically Porous Materials Based on Phase Separation (1991) . . . . . . . . . . . . . . . . . . . . . . . . Progress in Preparation of Aerogels (1992–1995) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Progress in Preparation of Mesoporous Materials (1994–1997) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Monolithic Columns for Chromatography and Elastic Gels (2000–2007) . . . . . . . . . . . . . . . . . . Aerogels of Various Compositions (Around 2010) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Concluding Remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

The history of sol‐gel technology has been reviewed on the basis of sol‐gelderived functional materials. When worldwide efforts for significant sol‐gel processing started to be made around 1970, it was regarded as a new method of preparing homogeneous glasses and ceramics at low temperatures, compared to the conventional melting and sintering method. After the sol‐gel preparation of inorganic–organic hybrid materials was proposed in 1984, however, glass and ceramic researchers started processing functional materials with optical, electrical, chemical, mechanical, or biomedical functions as well as advanced glasses and ceramics by sol‐gel method, and around 1995 sol‐gel technology attracted the people working in all the materials-related technology areas including electronics, chemistry, mechanics, pharmacy, and medicine. Since then, sol‐gel technology continued to progress by being applied to improvement of the performance of existing functional materials and exploration of novel functional materials. Detailed description of these situations will be made.

Introduction A routine sol‐gel processing for fabricating materials starts with a solution containing metal compounds as source of oxides. Metal compounds undergo hydrolysis and polycondensation near room temperature, solidifying into a wet gel. Vaporization of solvents near room temperature gives rise to an aimed gel product. In fact, there are many variations in starting material and processing condition, which give rise to a wide range of functional materials. Sol‐gel method is now regarded as one of the most powerful means for fabricating functional materials. With sol‐gel method, you may be able to improve the functional materials explored by other methods in the past or solve the problems regarding the fabrication process, as well as create novel functional materials. Furthermore, it is expected that sol‐gel method itself experiences much progress in the process of its application to fabrication of materials. Since sol‐gel method aims at fabricating some useful materials, applications are quite important. Bearing these ideas in mind, recent history of sol‐gel technology in terms of sol‐gel products will be described in this article. In the first section, history of overall sol‐gel products will be described, while

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in the following sections, four different types of sol‐gel products, that is, dye-sensitized solar cells, ferroelectric materials, biochemical and biomedical materials, and sol‐gel porous materials with controlled pores, are selected, and their history will be described separately.

History of Overall Sol‐Gel Materials Sol‐Gel Technology Before the 1960s Kistler (1931, 1932) reported preparation of aerogels. These aerogels may be regarded as the first significant sol‐gel products, although the word “sol‐gel” was not used. Kistler used water glass (aqueous solution of sodium silicate) as silica source and dried the resultant silica gel under methanol supercritical condition after exchanging water for methanol. Supercritical drying suppresses otherwise possible large shrinkage. Roy (1956) used metal alkoxide mixtures to construct phase equilibrium diagrams. Attention was paid to compositional homogeneity of gels prepared from alkoxide mixtures.

Sol‐Gel Technology in the Late 1960s to Early 1980s Around the year 1969 and later, the sol‐gel activity became high. Schröder (1969) deposited thin films on glass substrate from metal organic solution, and Mazdiyasni et al. (1969) prepared high-purity barium titanate powders from barium and titanium alkoxides. Roy (1969) prepared small silica glass pieces by heating gelled water glass. These works were interesting, but did not raise much attention among materials researchers. On the other hand, Dislich’s work (1971) on new routes to multicomponent oxide glasses published in 1971 attracted much attention of glass researchers. It was shown that gel powder derived from a homogeneous solution consisting of alkoxides of sodium, boron, aluminum, and silicon produced transparent Pyrex-type glass lenses on hot pressing at surprisingly low temperatures as low as 630  C. It is well known that Pyrex-type glass needs very high temperatures of more than 1500  C for fabrication by melt-quenching technique. Dislich’s work gave a big impact on glass researchers. Sakka et al. (1974) published a sol‐gel paper on preparation of TiO2–SiO2 amorphous materials at the 10th International Congress on Glass. Thereafter, sol‐gel research was directed toward preparation of bulk oxide materials. A special emphasis was laid on large pieces of silica glass (Yamane et al. 1979; Sakka and Kamiya 1979) and TiO2-doped SiO2 glass Sakka and Kamiya (1980), although bulk forms of transparent alumina were also prepared, and SiC ceramic fibers with very high strength (Yajima et al. 1976) attracted much attention. One of the reasons for focusing on bulk silica glass was to prepare silica preform for drawing optical fibers used in optical telecommunication which was emerging as technology of great importance, accomplishing very low optical loss of 0.2 dB/km around 1979.

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The First Sol‐Gel Workshop in 1981 In 1981 an international conference named “International Workshop on Glasses and Glass Ceramics from Gels” was held (Gottardi 1982) in Padova, Italy, which served in making the sol‐gel method familiar to the industries as well as academic world. In this workshop, there were 80 participants from eight countries including DDR, Egypt, England, France, Germany, Italy, Japan, and the USA. Eighteen research papers on sol‐gel science and technology were presented. After the International Workshop, the research on processing of bulk silica glass was actively continued (Zarzycky et al. 1982; Rabinovich et al. 1982; Wallace and Hench 1984; Klein et al. 1984; Adachi and Sakka 1987) until very large-size transparent silica glasses of, for example, 50 mm diameter and 2 m length (Toki et al. 1988) and large-size silica glass preform jackets for optical communication fibers (MacChesney et al. 1998) were obtained in 1988 and 1998, respectively. Besides bulk SiO2 glass, anti-reflecting coating of plate glass (Dislich and Hussmann 1981), sun-shielding coating of window glass (Dislich and Hinz 1982), In2O3–SnO2 conducting film on glass (Ogiwara and Kinugwa 1982), ZrO2–SiO2 fibers (Kamiya et al. 1980), Al2O3 fiber (Sumitomo Chemicals 1974), SiO2–Al2O3–B2O3 heat-resistant fiber (3M Company 1983), silica fiber (Sakka and Kamiya 1984), ferroelectric films (Wu et al. 1984), and so on were prepared in or by the first half of the 1980s. These sol‐gel products, except ferroelectric films, were almost completed and later improvement was continued. Later progress of ferroelectric materials is wonderful, which will be described in section “History of Sol‐Gel-Derived Ferroelectric Materials.”

Novel Materials Concept Derived from Sol‐Gel Technology Around 1984 In the second workshop held in 1983, Philip and Schmidt (1984) published a paper on preparation of inorganic–organic hybrid materials, which are transparent and homogeneous and can be used as hard contact lens. These materials were prepared by reactions in a mixture of epoxysilane, methacryloxysilane, alkoxysilane, titanium alkoxide, and monomeric methacrylate and are characterized by the presence of chemical bonds between organic species and inorganic ones in the matrix. In 1985, Schmidt (1985) gave a name ormosil or Ormocer. In 1984, another type of inorganic–organic hybrid materials was prepared by doping sol‐gel matrix with functional organic molecules. In fact, Avnir et al. (1984) prepared rhodamine 6G-doped silica gel, for instance, observing its spectral change due to trapping. Invention of these inorganic–organic hybrid materials provided the sol‐gel method with a big chance for fabricating innumerable functional materials by selecting compositions of inorganic–organic moieties and by using various functional organic molecules, becoming the basis for extensive development of novel materials together with the sol‐gel method so far explored. It is to be noticed that materials-minded researchers were aware of the versatility, flexibility, and possibility in tailoring the sol‐gel

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method. Certainly, the invention helped the sol‐gel method to be expanded to other areas than glass and ceramics, such as chemistry, electronics, architecture, mechanics, pharmacy, and medicine, slowly but steadily.

Development in 1985–1995: New Functional Materials by the Sol‐Gel Method Preparation of advanced materials in 1985–1995 was really remarkable as seen from the following examples. Photonic elements, such as phosphors (Reisfeld 1990a), dye lasers (Knobbe et al. 1990), nonlinear optical films (Zieba et al. 1992; Reisfeld 1990b), photochemical hole-burning memories (Tani et al. 1985), photochromic molecules (Levy et al. 1989), photochemical sensors (Samuel et al. 1994; MacCraith et al. 1994), reflectors for head-up display of cars (Hattori et al. 1989), and so on were prepared by sol‐gel method. Sol‐gel processing was also applied to electronic materials, such as transparent conducting films (Kodaira et al. 1990; Furusaki et al. 1994) used as electrode for dye-sensitized solar cells (O’Regan and Grätzel 1991; Grätzel 1994) (refer to section “History of Dye-Sensitized Solar Cells”), electrochromic cells, and so on. It is commented that the electrical conductivity of the sol‐gel ITO films obtained in the above researches would be the highest among the sol‐gel ITO films and that usually the conductivity of ITO films obtained by sputtering is higher than that of the sol‐gelderived ones. Soon after the discovery by Bednorz and Mueller (1986), sol‐gel processing was applied to processing of superconducting oxide films and fibers of the Y–Ba–Cu–O system (Nasu et al. 1987; Monde et al. 1988) and the Bi–Sr–Ca–Cu–O system (Agostinelli et al. 1988; Tohge et al. 1990; Zhuang et al. 1990). Research on ferroelectric materials showed a considerable progress (Uhlmann et al. 1992), which is referred to section “History of Sol‐Gel-Derived Ferroelectric Materials.” Chemical and mechanical protection of substrate was studied with the following titles: passivation of electrical circuit by hybrid film (Popali et al. 1994), strengthening of SiO2 glass by silica film (Fabes et al. 1986), improvement of acid durability of iron and steel plate by coating with CH3Si(OC2H5)3 solution (Murakami et al. 1989), and providing with water repellency by ZrO2 or SiO2–ZrO2 coating (Tohge et al. 1987). Micropatterning on SiO2–TiO2–polyethylene glycol film was also carried out by stamping (Tohge et al. 1988). The proposal of Philip and Schmidt (1984) for inorganic–organic hybrid materials opened the door for coating plastic substrates with abrasion-resistant materials. The sol‐gel method is quite suitable for this type of coating, because the processing can be made at low temperatures less than 180–200 C, where the plastic substrates are not decomposed and they are harder than polymers and less brittle compared to inorganic glasses and ceramics. Thus, inorganic–organic hybrid coatings were soon applied to ophthalmic lenses (Schmidt et al. 1988), and similar hybrid coatings have been commercially available since then (Schottner et al. 2003).

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The concept of inorganic–organic hybrid materials created in 1984 expanded the area of the sol‐gel method even to biochemical and biomedical technologies. As to this matter, refer to section “History of Sol‐Gel Biochemical and Biomedical Materials.”

Progress of Sol‐Gel Technology in 1995–2005 One of the most important topics in the area of sol‐gel optical materials in this decade is sol‐gel preparation of photocatalyst. The concept of photocatalyst emerged from the discovery of Honda–Fujishima effect (Fujishima and Honda 1972), in which water is decomposed into H+ and OH ions at the surface of TiO2 crystal on UV light exposure. Hashimoto and Fujishima (1994) showed that TiO2 photocatalysts decompose contaminants, mainly consisting of organic compounds and bacteria, cleaning the surface of walls, windows, and so on (Ikeda et al. 1997; Ishibashi et al. 1998). Hydrophilicity of the surface of TiO2 crystals caused by UV exposure promotes elimination of contaminants (Nakajima et al. 2001). Around 2000, sol‐gel-prepared photocatalyst was commercially applied to windows of buildings as self-cleaning coating (Tanaka and Obana 2004). In the 2000s, visible light-responsive photocatalysts, which work within buildings and houses, were explored by doping TiO2 with nitrogen (Asahi et al. 2001) and both Sb and Cr (Kato and Kudo 2002) and Cr or V (Anpo and Takeuti 2003). Titanium oxide is the most important compound as photocatalyst for the reasons of high photocatalytic performance, good chemical stability, high durability to UV and visible light exposure, and low cost of raw material. The photocatalytic performance depends not only on the microstructure of the film including porosity and pore size but also composite formation or doping. The effects of dopants, such as CaO, MgO, CaCO3, Al2O3 and Fe2O3 (Ichiura et al. 2003), WO3 (Miyauchi et al. 2002), and so on were examined for UV-effective TiO2 photocatalysts. Such efforts have been continued until present. In the area of electronic materials, ferroelectric materials experienced marked expansion and regulation on the use of lead. For details, refer to section “History of Sol‐Gel-Derived Ferroelectric Materials.” A certain progress was seen in sol‐gel coating of plastic substrates in this decade. Antiglare conducting ITO coating was prepared on plastic substrates by using inorganic–organic hybrid sols containing the highest possible amount of already conducting crystalline In2O3–SnO2 nanoparticles (Aegerter and Al-Dahoudi 2003). The layers were cured by UV light irradiation of polymerizable organic additives. Asakuma et al. (2003) obtained conductive, transparent In2O3–SnO2 coatings on polyethylene terephthalate and polyimide sheets via a sol‐gel route assisted with a UV laser beam. The film was crystallized on exposure to a low-fluence UV beam. For the progress of biochemical and biomedical technology, refer to section “History of Sol‐Gel Biochemical and Biomedical Materials.”

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Progress of Sol‐Gel Technology in 2005– 2015 Also in this decade, the sol‐gel technology continued to progress. One of the most well-known topics in the area of optics is popularization of white light-emitting diode (LED) for lighting or illumination. White LED was made possible by invention of blue ray-emitting LED (Nakamura 1991). Commercialization of white LED lighting system rapidly progressed. Sol‐gel technology is very important for tailoring phosphors used for white LED. It is known that the highest color rendering white LED consists of a blue LED and a blue light-excited yellow phosphor of Ce/YAG (yttrium aluminum garnet). In order to increase heat resistance of phosphors, ceramics of Ce/YAG or Ce/GdYAG are used (Nishiura et al. 2011a, b). Also, for higher color rendering white LED, RGB-emitting Tm3+-, Tb3+-, and Eu3+-doped CaW1-xMoxO4 phosphors are proposed, and wavelength tunability between cool, natural, and warm white lights is investigated (Mickens et al. 2012). It is to be noted that sol‐gel-derived inorganic–organic hybrids (Kim et al. 2010b) and silica (Wang and Huang 2012) are proposed as materials for encapsulants of phosphors. In this period, efforts were made in improving the performance of photocatalysts or in finding novel photocatalysts. Katsumata et al. (2010) showed that a new photocatalyst of niobium oxide applied on soda lime glass reacted with sodium ions diffusing out of the glass substrate, being stabilized as NaNbO3 when heated at temperatures higher than 450  C. The NaNbO3 film gave rise to excellent selfcleaning property for glass under UV irradiation, although its photocatalytic oxidation activity was poor. Katsumata suggested that this film is more suitable as selfcleaning agent for vehicle than TiO2 photocatalyst, because the film is firmly attached to vehicle windows by NaNbO3. Wu et al. (2012) proposed another new photocatalyst of nanosized ZnNb2O6 as exhibiting very high photocatalytic activity, due to its favorable crystallinity and surface area. One of the developments in the area of electronic materials in this decade is the application of sol‐gel method to preparation of novel transparent oxide conductors discovered before and fabricated by sputtering and other physical techniques. Sol‐ gel preparation of amorphous In–Ga–Zn–O, which was regarded as the most potential candidate for the transparent conductive oxides by H. Hosono, has been carried out (Lin et al. 2014). The effect of composition of films on the film structure was precisely examined by adopting sol‐gel processing technique, and it was found, for instance, that an increase in Ga content makes preparation of amorphous films easy. Heating at temperatures lower than 300  C could produce good In–Ga–Zn–O films. CuAlO2 film is an important p-type transparent conductor oxide discovered by Kawazoe et al. (1997). As to this polycrystalline compound, the presence of impurity crystals CuO and CuAl2O4 in synthesized materials was a problem to be solved. Li et al. (2010) applied the sol‐gel method to processing of this film and prepared pure thin films by annealing the solution-deposited film on sapphire or quartz glass at 850  C in nitrogen flow under atmospheric pressure. The film showed low room temperature resistivity and high optical transmission in visible region. Another interesting transparent conductive film is Nb-doped TiO2 (TNO) polycrystalline

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film. So far, ITO films have been widely used as electrode materials, due to the low resistivity reaching 1  10 4Ωcm. However, indium is a rare element and TNO is regarded as a substitute replacing ITO. Zhao et al. (2010) prepared TNO films by sol‐ gel processing, obtaining resistivity of 19.3 Ωcm. This indicates that further research is needed to obtain meaningful results for TNO. As to development of solar cells and ferroelectric materials, refer to sections “History of Dye-Sensitized Solar Cells” and “History of Sol‐Gel-Derived Ferroelectric Materials,” respectively. The sol‐gel technology has been receiving considerable attention in processing of lithium-ion batteries. They are currently the most popular batteries for portable electronic devices, electric vehicles, and airplanes because of their high energy density, good cycling stability, and high voltage of 3.9 V. Most of lithium-ion batteries use LiPF6 as electrolyte, LiCoO2 or LiMn2O4 as cathode material, and carbon as anode. Research works are mainly focusing on cathode materials, for example, the use of LiCoPO4 (Vasanti et al. 2008; Vadivel et al. 2009), modification of LiCoPO4 by sol‐gel method (Rajalaksh et al. 2013; Sarapulova et al. 2012), and use of sol‐gel-derived LiNi1/3Co1/3Mn1/3O2 materials (Liu et al. 2012) and Li4Ti5O12 (Mosa et al. 2014) are proposed. Al-doped Li7La3Zr2O12 was prepared as electrolyte (Takano et al. 2014). Besides these, all-solid-state lithium-ion batteries are prepared (Van den Ham et al. 2015; Ohtomo et al. 2013). It is noticed that sol‐gel method is valuable for studies on the effect of doping or composition also in lithium-ion batteries. Considerable progress has been made in the area of sol‐gel coating of plastics in this decade. Tanaka et al. (2011) explored superdurable hard coat film based on fluorinated polyhedral oligomeric silsesquioxane-terminated polymer (Koh et al. 2005). Dinelli et al. (2011) used microwaves to enhance the scratch resistance of coating on polycarbonate. Kajiwara et al. (2009) also used microwave-assisted sol‐gel method to prepare intense luminescent hybrids with high photostability and low leachability. A novel and versatile method of applying large area ceramic thin films on plastics was proposed by Kozuka et al. (2012a, b, 2013) recently. In this method, a gel film is deposited on a silicon substrate coated beforehand with a release layer, the gel film is fired into a ceramic film, and, then, it is transferred onto plastics by softening the plastic surface. The advantage of this method is the possibility of fully crystallizing functional ceramic films beforehand and versatility expressed by being applicable regardless of the combination of ceramics and plastics. Patterned films can be also prepared by this technique. For progress in biological and biomedical materials, refer to section “History of Sol‐Gel Biochemical and Biomedical Materials.” So far, the history of overall sol‐gel technology was described in terms of sol‐gelderived functional materials. It was shown that sol‐gel technology continued to progress during these several decades. The following sections will be devoted to historical description of four kinds of materials, that is, dye-sensitized solar cells, ferroelectric materials, biochemical and biomedical materials, and porous materials with controlled pores, separately.

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History of Dye-Sensitized Solar Cells Silicon-based solar cells are now widely prevailing as one of the sustainable energy sources. The role of sol‐gel method in such silicon solar cells is limited to coatings on cover glass, such as silica coating film for protection of silicon and anti-reflecting film on cover glass for avoiding loss of light, because silicon films and transparent conductor oxide films (ITO) as electrode are fabricated mainly by sputtering. In this section, history of TiO2-based dye-sensitized solar cells will be considered.

Grätzel’s Dye-Sensitized Solar Cells In 1991, O’Regan and Grätzel (1991) reported on a novel photoelectrochemical solar cell, called dye-sensitized solar cell, which converts lights to electricity via mechanism (Grätzel 1994) different from the abovementioned solid-state pn-junction silicon solar cell. The novel dye-sensitized solar cell is made up of [sheet glass/ FTO conductor/liquid electrolyte/dye layer (sensitizer)/TiO2 film/FTO conductor/ sheet glass]. In this solar cell, dye molecules absorb visible lights and produce excited electrons, which are inserted into the conduction band of TiO2 film. These electrons move to the outer circuit, working as electricity. The electrons come back to the TiO2 film through the electrolyte, which contains redox pairs of iodine I and iodide anion I . It is easily understood that in these cells, the larger is the surface area of the dye layer and, accordingly, that of TiO2 film, the better is the cell performance. With sol‐gel technique, Grätzel prepared highly porous TiO2 films of 10 μ thickness, in which the surface area of TiO2 layer is several hundreds cm2 for 1 cm2 of illuminated geometrical area. The high conversion efficiency of about 8% reported by Grätzel (1994) startled those who were interested in solar cells. Afterward, research works seeking for more efficient TiO2 layer of large surface area and dye molecules absorbing more solar energy and bonding with TiO2 more tightly have been continued by many people, but it was not easy to exceed 8% in conversion efficiency.

Dye-Sensitized Solar Cell with Solid Electrolyte In the dye-sensitized solar cell invented by Grätzel, liquid electrolyte has been employed. As a method of assuring long-term stability of the cell, solid-state dye-sensitized solar cell was sought for. Actually, Meng et al. (2003) prepared a dye-sensitized solar cell-containing CuI, a hole conductor, instead of the liquid electrolyte, obtaining conversion efficiency of 3.8%. Another problem in liquid electrolyte may consist in degradation of the cell performance due to the vaporization loss of the electrolyte solvent. Lan et al. (2010) prepared dye-sensitized solar cell by using stable organic–inorganic composite electrolyte which slowly solidifies as a result of reaction between

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components. The energy conversion efficiency was around 3.5%. Improvement of the efficiency is needed.

Perovskite Solar Cell In 2009, a novel dye-sensitized solar cell, later called perovskite solar cell, was explored by Miyasaka’s group (Kojima et al. 2009), who replaced the organic dye of the dye-sensitized solar cell by organic–inorganic perovskite crystals of the composition CH3NH3PbI3. Then, the iodine-containing liquid electrolyte used in the above dye-sensitized solar cell was also replaced by a solid transporting positive holes. The perovskite solar cell thus improved showed the efficiency of 10.9% (Lee et al. 2012). The perovskite solar cell is made up of [Au/Spiro-OMeTAD/CH3NH2PbI3/TiO2/ FTO glass]. The organic–inorganic hybrid perovskite crystalline layer of 300 nm thickness absorbs all visible lights, converting them to electrons with a quantum efficiency of unity. The electrons are transported to cathode through TiO2 film, while holes are transported to Au electrode through Spiro-OMeTAD. After the invention of the perovskite solar cell, a great many research reports showing higher conversion efficiencies have been published. For example, a report (Jeon et al. 2015) indicated a very high efficiency reaching 18%. Commercialization in the future may be expected for the perovskite solar cells.

History of Sol‐Gel-Derived Ferroelectric Materials Ferroelectric materials show hysteresis loops in polarization (P) vs. electric field (E) curve and are characterized by high dielectric constant and large piezoelectric, pyroelectric, and electrooptic effects, although there are many variations. On the basis of these unique properties, ferroelectric materials have been applied to fabrication of capacitors, oscillators, actuators, temperature sensors, and nonvolatile random access memories. Generally, we apply BaTiO3 for capacitors, Pb(Ti,Zr)O3 and LiNbO3 for piezoelectric devices, and LiTaO3 for pyroelectric devices. More specifically, SrBi2(Ta,Nb)2O8 is applied for nonvolatile ferroelectric random access memories (NV-FeRAM). It should be remembered that ferroelectric materials are often used in the form of thin coating films, which can be fabricated by sol‐gel method. Bearing these in mind, history of ferroelectric materials will be described.

Sol‐Gel Preparation of Ferroelectric Materials Around 1970, Mazdiyasni’s group prepared dense ceramic of BaTiO3 (Mazdiyasni et al. 1969) and transparent ceramic of lead lanthanum zirconate titanate (PLZT) (Brown and Mazdiyasni 1972) by sintering powder mixtures derived from solution of metal alkoxides. This sol‐gel method starting from solutions made it possible to obtain dense and homogeneous products by heating at temperatures 300–400  C

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lower than the temperatures employed in the conventional method starting from a mixture of solid powders. However, these works did not cause an extensive development of sol‐gel preparation of ferroelectric materials before the 1980s.

Sol‐Gel Preparation of Ferroelectric Materials in 1980–1995 In the early 1980s, sol‐gel preparation of ferroelectric materials started with the search for possibility of forming thin films without cracks (Sakka and Kokubo 1983) and obtaining aimed ferroelectric crystals (Riman et al. 1984). Progress of sol‐gel processing of ferroelectric materials was slow in the middle 1980s, but a considerably large number of papers were published after 1987. Hirano and Kato (1988a, b; 1989) fabricated LiNbO3 thin films on silicon and sapphire crystals for the purpose of obtaining high-reflective index electrooptic materials. By using double-metallic alkoxide solution, the film could be crystallized with preferred orientation at low temperatures as low as 250  C. Incidentally, the use of double alkoxides serves in avoiding formation of non-ferroelectric pyrochlore otherwise accompanied by crystallization of perovskite. Double metallic alkoxide precursors are tabulated by Kessler (2004). At the end of the 1980s, certain researchers started measuring ferroelectric properties of sol‐gel-prepared materials. Table 1 shows values of remnant polarization Pr and coercive field Ec of ferroelectric thin films obtained by Xu et al. (1990) and Scott et al. (1991). It is seen from Table 1 that pretty good ferroelectric thin films are obtained. Uhlmann et al. (1992), reviewing papers on sol‐gel preparation of ferroelectric materials published so far, expressed the opinion that the sol‐gel method might be more effective than sputtering method in developing new compositions and that for the application stability of properties in terms of switching cycles is of primary importance.

Progress in 1995–2005 and Tasks Imposed by Restriction on Lead-Containing Materials After the middle of the 1990s, sol‐gel processing of ferroelectric materials experienced quite marked expansion and progress as can be seen in the paper published by Uhlmann et al. (2000). Large lists given in this paper tell that the progress can be Table 1 Properties of ferroelectric thin films obtained around 1990 BaTiO3 Pb(Zr0.5Ti0.5)O3 0.15%Nb2O5 doped Pb(Zr0.5Ti0.5)O3 Sr0.6Ba0.4Nb2O6 KNbO3

Pr (μC/cm2) 21 32

Ec (KV/cm) 83 15.7

Thickness (μ) 0.8 0.8

Reference Xu et al. 1990 Xu et al. 1990

15 34 13

45 51 86

0.22 0.9 0.7

Scott et al. 1991 Xu et al. 1990 Xu et al. 1990

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Table 2 Ferroelectric properties of SrBi2Ta2O9 thin films measured by different groups Composition SrBi2Ta2O9

Pr (μC/cm2) 10

Ec (KV/cm) 38

Stability 109

Thickness (μ) 0.28

SrBi2Ta2O9 SrBi2Ta2O9 Sr0.7Bi2.3Ta2O9+α

5 7.5 7

55 45 50

1011 4  109 1010

0.2 0.5 0.8

Reference Amanuma et al. 1995 Chu et al. 1996 Boyle et al. 1996 Koiwa et al. 1996

Stability: number of repetition cycles until occurrence of fatigue

found in all kinds of ferroelectric materials such as nonvolatile ferroelectric random access memories, piezoelectric materials, and pyroelectric materials. These expansion and progress are assumed to be caused by the fact that researchers in the field of electronics and mechanics as well as glass and ceramics researchers joined the research of ferroelectric materials using sol‐gel processing. Table 2 compares the results of measurements by four research groups (Amanuma et al. 1995; Chu et al. 1996; Boyle et al. 1996; Koiwa et al. 1996). It is seen from Table 2 that research was carried out by a large number of researchers and good and similar values were obtained for remnant polarization and coercive field. Furthermore, the large numbers of repetition cycles before fatigue indicate that excellent ferroelectric films can be fabricated by sol‐gel processing. In 2003, when the sol‐gel-derived ferroelectric materials were showing high performance, the Restriction of Hazardous Substances (RoHS) Directive has been in force. The directive states that electrical and electronic equipment sold after July 2006 must not contain hazardous substances such as Hg, Cd, and Pb. This news was shocking to the researchers of ferroelectric materials, because lead-containing Pb(Zr, Ti)O3 and Pb(Mg,Nb)O3 were regarded as the most important, representative ferroelectric materials. However, some of the researchers predicted the restriction on the use of lead beforehand, aiming at improvement in performance of existing leadless ferroelectric materials or development of new leadless ones. As examples of the latter case, bismuth system layer-structured ferroelectric coating films of the composition CaBi4Ti4O5 were developed by sol‐gel method (Kato et al. 2001, 2004). The resultant film showed excellent ferroelectric properties and high stability: remnant polarization Pr = 25 μC/cm2 and coercive field Ec = 306 KV/cm. The film also showed a large piezoelectric coefficient d33 = 30 pm/V. These values indicate that sol‐gel-derived leadless ferroelectric materials can replace conventional lead-containing piezoelectric materials exemplified by Pb(Zr,Ti)O3.

Progress in 2005–2015: Sol‐Gel Processing of Multiferroic Materials Multiferroic materials have been intensively studied in this period (2005–2015). In these materials ferroelectricity and ferromagnetism coexist. The coupling between magnetic and electrical ordering may lead to the magnetoelectric effect in which

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magnetization is a function of electric field and dielectric polarization is a function of magnetic field. Multiferroic materials have attracted much attention of solid-state physicists as a fundamental physics problem and technologists as materials creating applications for novel memory devices and piezoelectrics. For fundamental and application research, many researchers took up BiFeO3 which is ferroelectric with Curie temperature Tc of about 810  C and antiferroelectric with Neel temperature of about 380  C. Before 2000 multiferroic materials were prepared by sputtering. In the early 2000s, application of sol‐gel method for preparation of BiFeO3 film started. Later, Habouti et al. (2007) and Liu and Wang (2008) indicated that simple BiFeO3 composition does not easily produce high-performance coating films, due to the appearance of the second phase, high leaky current, and so on. In order to improve the performance, doping BiFeO3 with Ti, Mn, and Ce and forming BiFeO3-based solid solution were tried. For instance, Sakamoto et al. (2006, 2009) prepared BiFeO3–PbTiO3 solid solution thin films and Mn-doped solid solution. It can be seen that the latter thin films showed relatively saturated ferroelectric hysteresis loop in polarization versus electric field curve at room temperature compared with the former. Both the films showed distinct ferromagnetic hysteresis in the magnetization vs. magnetic field curve. Aside from the BiFeO3-based multiferroic materials which show both ferroelectricity and ferromagnetism, multiferroic materials which show a considerable magnetoelectric coupling coefficient have been investigated. Chen et al. (2013) showed that composite films consisting of Bi3.25Nd0.75Ti3O12 and La0.6Ca0.4MnO3 exhibit a large magnetoelectric effect as well as both good ferroelectric and magnetic properties.

Recent State of the Art of Ferroelectric Materials The present task concerning ferroelectric coating film is to improve the ferroelectric, piezoelectric, and pyroelectric properties and to increase the long-term stability of the property, as it was in past decades. In order to achieve this, the effect of large and small changes of the compositions, constitution, particle size of the film, and processing conditions has to be controlled. Leadless BiFeO3-based ferroelectric materials are candidates for highly ferroelectric ones. As to the composition, sodium-containing compositions have been recently proposed (Kim et al. 2010a; Fang et al. 2011; Zhu et al. 2015). In this case, the effect of sodium on the chemical stability of the materials has to be checked. The sol‐gel method is quite suitable for carrying out such tasks, because precise and detailed modification in composition and processing is possible in the sol‐gel method.

History of Sol‐Gel Biochemical and Biomedical Materials Generally, biochemical and biomedical materials should be nontoxic and biocompatible for the living body to be protected. Besides, in sol‐gel processing of biochemical and, especially, biomedical materials, high concentrations of alcohols and

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acids or alkalis have to be avoided, although in conventional sol‐gel synthesis, they are often employed as solvent and catalyst, respectively.

Hench’s Suggestion on Compatibility of Silicates The concept of inorganic–organic hybrid materials was proposed in 1984 (Philip and Schmidt 1984; Avnir et al. 1984). Soon after that time, Hench, who invented biologically active, artificial bone glass of the system Na2O–CaO–P2O5–SiO2 (Hench 1974), suggested (Hench and Orcel 1986) that certain silicates are compatible with living organism and tissues. It is assumed that both these concepts became a trigger to application of the sol‐gel-derived gels for encapsulating biochemical and biomedical molecules and living tissues. It is noted also that the sol‐gel method might be suited for these applications, due to the fact that gels are characterized by continuous pores and controllability of pore size and pore distribution.

Starting and Rapid Development of Sol‐Gel Preparation of Biochemical and Biomedical Materials in the 1990s Enzymes (proteins) were encapsulated in sol‐gel-derived gels in 1990 (Braun et al. 1992; Ellerby et al. 1992). A few years later, a great many papers on sol‐gelderived materials with biochemical and biomedical functions were published. Some examples are shown in Table 3. Biosynthesis takes place in gels in which enzymes or enzyme-producing microorganisms and cells are immobilized (Braun et al. 1992; Ellerby et al. 1992; Pope 1995a). Biosensor consists of gel-entrapped molecules which react with molecules to be analyzed, resulting in a change in spectra of optical absorption or fluorescence (Braun et al. 1992). The application of sol‐gel method to implant consists in incorporating living cells, tissues, and organs into gels and the gels into animals (Pope 1995b). Medical diagnosis by immunoassay utilizes antibody–antigen reactions in porous gel matrix entrapping specific protein (Livage et al. 1996; Turniansky et al. 1996). Probably, all sol‐gel applications shown in Table 3 might be surprising for most of the investigators and engineers in the field of ceramics and glasses. Among them, sol‐gel encapsulation of living tissues (pancreatic cells) by silicate gel reported by Pope et al. (1997) would be more striking.

Progress in 2000–2006 In the period of 2000–2006, a certain progress was seen in encapsulation of pancreatic cells. Sakai et al. (2002a, b, 2003) prepared sol‐gel-derived microencapsulation which protects the encapsulated Langerhans islets from the immune antibodies and at the same time releases insulin secreted by the cells. Alginate/amino-

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Table 3 Sol-gel processed biochemical and biomedical materials Application Bioreactor element (biosynthesis)

Biosensor (fluorescence, color) Implant of living tissues and organs

Medical diagnosis and test

Entrapped materials Enzyme; lipase Hemoglobin Microorganism Plant cell Animal cell

Entrapping materials Silica gel Silica Silica gel

Phycobiliprotein Glucose oxidase Hepatocyte tissue Pancreatic islet Resorbable artificial bone Antibodies Antibody–protein antigen Anti-atrazine antibodies

Gel Gel Silica gel Silica gel

Silica gel

Reference Braun 1990; Reetz et al. 1996 Ellerby et al. 1992 Pope 1995a Campostrini et al. 1990 Boninsegna et al. 2012 Chen et al. 1996 Sampath et al. 1996 Pope 1995b Pope et al. 1997 Gerber et al. 2003 Liverge et al. 1996 Roux et al. 1997 Turniansky et al. 1996

propyl-silicate/alginate microcapsules encapsulating islet cells were implanted into diabetic mice. This reversed the diabetes in animals within 1 day after implantation, and the normoglycemic state was maintained for 2–3 months. Another big topic in these years is the use of sol‐gel hybrid films for lab-on-a-chip devices (Durucan and Pantano 2004). Lab-on-a-chip devices are microwell plates or DNA arrays, which carry a very large number of finely divided test zones, so that various information on chemical performances, reaction between proteins and DNAs, and so on are obtained in a short time. Usually, plate glasses with sol‐gelderived inorganic–organic hybrid coating films are used as base plate. Other sol‐gel biochemical and biomedical technologies developed in these years are the proposal of the use of sol‐gel-derived inorganic–organic hybrids as carrier for drug delivery (Boettcher et al. 1998); studies on hybrid systems based on chitosan (Yokogawa et al. 2001; Shirosaki et al. 2005), which is found in nature as shells of crabs and lobsters and is attracting much attention because of its high biocompatibility; and the use of poly(lactic acid)–calcium carbonate composites as resorbable scaffold in bone implantation (Maeda et al. 2006).

Progress in 2006–2015 The latest decade (2006–2015) can be defined as the period for improvement or extension of biochemical and biomedical materials, especially encapsulation or immobilization by inorganic–organic hybrids. Encapsulation of living yeast (Guan et al. 2008), living bacteria and algae (Soltmann and Böttcher 2008), and penicillin G acylase (Bernardino et al. 2011) was reported, for instance.

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The advantage of porous nature of sol‐gel materials was reconfirmed with bioactive scaffold (Almeida et al. 2011). It was shown in this paper that sol‐gelderived SiO2–CaO–P2O5 porous glass with a dual pore structure consisting of interconnected pores of both 100s of micrometers and 10s of nanometers in size has good potential as a scaffold in bone tissue application. In drug delivery systems, drugs entrapped in nanoparticle matrices are carried onto the particular localized area of tumor, for instance (Sasaki and Akiyoshi (2012)). For hyperthermia, Azevedo et al. (2014) prepared the composite of magnetite nanoparticles coated by mesoporous silica and a temperature-responsive poly (N-isopropylacrylamide) polymer, which makes it possible to target localized heating in vivo by an alternating current magnetic field.

History of Processing Porous Materials with Controlled Pores Sol‐gel-derived inorganic gels are usually accompanied by fine pores, in contrast to organic gels which usually have no explicit pores except for special case of, for example, resorcinol-derived gel (Fricke and Gross 1994), as far as I know. Examples of three kinds of SiO2 xerogels with uncontrolled pores are shown (Sakka et al. 1993), in which the average size of pores and the porosity are 2.5 nm and 46%, 16 nm and 73%, and 5 μ and 60%, respectively. Xerogel means a gel which was dried by vaporization at ambient pressure. In this section, history of sol‐gelderived porous materials with controlled porosity, pore size, and/or pore distribution will be described. Such gels are classified into three types, according to difference in processing: “aerogels” processed by supercritical drying, “mesoporous materials” synthesized with templates, and “gels with hierarchical pore distribution” formed by concurrent phase separation and gelling. Aerogels are taken up as porous materials with controlled porosity.

Aerogels (1931, 1968) Kistler (1931) reported preparation of silica aerogels using sol‐gel method combined with supercritical drying. Kistler used water glass as starting material. After exchanging water of gel with alcohol, the gel was subjected to alcohol supercritical drying. This is the first significant sol‐gel work, although there was no nomenclature “sol‐gel” at that time. Silica aerogel is a highly porous solid and is characterized by low density, low refractive index, and very low thermal conductivity. In spite of those unique properties, SiO2 aerogel was hard to be utilized because of its mechanical fragility and weakness to moisture. Kistler prepared also aerogels of alumina, iron oxide, tin oxide, and so on. In 1968, Nicolan and Techener (1968) proposed the use of alkoxysilanes as starting materials for preparing silica aerogels. A starting solution of the tetramethoxysilane–alcohol–water system was subjected to hydrolysis–polycondensation, and the resultant humid gel was dried under the

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supercritical condition of methanol, which gave rise to transparent aerogels of very large size, for example, 50  60  80 cm, without fracture and shrinkage.

Mesoporous Materials Based on Self-Assembly (1990, 1992) Around 1990, preparation of both “mesoporous materials” and “hierarchical porestructured gels” started. Yanagisawa et al. (1990a, b) prepared novel ordered mesoporous silica using the product of reaction of layered silicates with organoammonium surfactants. In 1992, Kresge et al. (1992) and Beck et al. (1992) reported the work in which novel ordered silica thin films were prepared by calcining ordered mesostructured thin organic-silica hybrids fabricated from the silicate solution containing surfactants or liquid crystals as templates. The thin films contain uniform pores arranged in an ordered array as found in crystalline materials, such as zeolite. The pore size changes from 2 to 20 nm. Those papers gave rise to extensive development of templated synthesis of nanostructured mesoporous materials.

Hierarchically Porous Materials Based on Phase Separation (1991) In 1991, Nakanishi and Soga (1991) and Nakanishi et al. (1991) reported on the novel method of preparing silica with hierarchical pores. Gels having both macropores and mesopores with controlled pore size and pore size distribution were prepared through sol‐gel reaction accompanied by phase separation in the starting solution containing water-soluble polymers and alkoxysilanes. Typical resultant silica gels contained micrometer-size frameworks with nanometer-size continuous pores (0.5–5 μ), and silica frame works with nanometer-size (5–50 nm) open pores.

Progress in Preparation of Aerogels (1992–1995) In 1992, a novel method for aerogel formation, that is, solvent exchange method, was proposed by Tillotson et al. (1992). This method uses drying under the CO2 supercritical condition (Tc, 31.1  C; Pc, 73.9 bar) which is much milder than the methanol supercritical condition (Tc, 239.4  C; Pc, 80.9 bar). The use of CO2 supercritical condition makes it possible to carry out supercritical drying at low temperatures close to ambient temperature, although exchange of alcohol by CO2 before drying is time-consuming. In 1994 another interesting proposal for facilitating production of aerogel-like xerogel was made. Ambient pressure drying for synthesizing very low-density gels was first attempted by Einarsrud and Haereid (1994). This was accomplished by strengthening the gel structure in order to suppress possible large shrinkage before ambient pressure drying. Later, such very low-density xerogels were called aerogel. Since then, a large number of papers have been published on this subject (Aravind et al. 2010).

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In 1995, Yokogawa and Yokoyama (1995) prepared hydrophobic silica aerogel. Hydrophobicity was given by trimethylsilylation with hexamethyldisilazan, which reacts with hydroxy groups on the surface of silica, capping those groups. This work contributed to the enlarged application of silica aerogels including commercialization.

Progress in Preparation of Mesoporous Materials (1994–1997) As to mesoporous materials, solvent evaporation method was proposed and reported (Ogawa 1994). That is, the evaporation-induced self-assembly was utilized to prepare ordered mesoporous films by using the surfactant and the inorganic precursor dissolved in alcohol. During dip or spin coating, the concentration of surfactant drastically increases, which causes self-assembly. Later, in 1997 Lu et al. (1997) reported the continuous formation of ordered mesoporous thin coating films. Since then various kinds of mesoporous materials have been developed.

Monolithic Columns for Chromatography and Elastic Gels (2000–2007) Since the first paper (Nakanishi and Soga 1991) was published in 1991, intensive research works have been carried out on “gels with hierarchical pores” as doublepore system gels. The effects of processing condition and composition of starting solution on the size and volume fraction of macropores and mesopores and pore surface chemistry have been investigated. With those researches, monolithic gels were applied as monolithic columns for high-performance liquid chromatography (Nakanishi et al. 2000; Nakanishi 2004) around the year 2000. Kanamori et al. (2007, 2008) reported preparation of novel aerogels with unusual mechanical strength against uniaxial compression. The novel aerogels having methylsilsesquioxane networks were prepared from CH3Si(OCH3)3 (MTMS) solution containing surfactant and urea as phase separation-inducing agent and as accelerator for the condensation reaction, respectively. The optimized aerogels dried under a supercritical condition showed not only similar properties as conventional silica aerogels such as high transparency and high porosity but also outstanding mechanical strength against compression. The compression test showed that an MTMS aerogel sample prepared from starting solution consisting of water–acetic acid–MTMS–CTAB (cetyltrimethylammonium bromide) shrinks up to 84% in linear scale upon uniaxial loading and springs back to the original size upon unloading.

Aerogels of Various Compositions (Around 2010) So far, many works on porous silica have been introduced. It should be known, however, that many kinds of porous materials other than silica have been prepared by sol‐gel method. Some of the examples of aerogels are shown: Al2O3 aerogel

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(Bono et al. 2010), TiO2–SiO2 aerogel (Cao et al. 2008), Fe2ZnO4 aerogel (Brown and Hope-Weeks 2009), high-strength vanadia (VOx) aerogel (Luo et al. 2008), CdS aerogel (Mahanan and Brock 2006), carbon aerogel (Lorjai et al. 2009), and so on.

Concluding Remarks Although Kistler’s paper on preparation of aerogels published in 1931 is recognized as the first significant sol‐gel work, it is after the late 1960s that many materials researchers started having interest in the sol‐gel method as novel method for processing glass and ceramics. They applied the sol‐gel method for fabricating bulk silica glass, functional ceramics, and coating films. In 1984, a big chance for expansion was brought about to the sol‐gel method by the invention of inorganic–organic hybrid materials by the two groups of Schmidt and Avnir. Due to this invention, the sol‐gel method spreads to other areas than glass and ceramics, that is, to areas of chemistry, electronics, mechanics, architecture, pharmacy, and medicine. In 1995–2005, therefore, the sol‐gel method was applied to preparation of all kinds of older and new functional materials and to improvement of quality of functional materials. This progressive tendency of sol‐gel method continued thereafter until the present flourishing state we see now.

References Adachi T, Sakka S. Preparation of monolithic silica gel and glass by the method using N,N-dimethylformamide. J Mater Sci. 1987;22:4407–10. Aegerter MA, Al-Dahoudi N. Wet –chemical processing of transparent, and antiglare conducting ITO coating on plastic substrates. J Sol–gel Sci Technol. 2003;27:81–9. Agostinelli JA, Paz-Pujalt GR, Mehrohtra RC. Superconducting thin films in the Bi-Sr-Ca-Cu-O system by the decomposition of metallo-organic precursors. Physica C. 1988;156:208–12. Almeida RM, Gama A, Vueva Y. Bioactive sol–gel scaffolds with dual porosity for tissue engineering. J Sol–gel Sci Technol. 2011;57:336–42. Amanuma K, Hase T, Miyasaka Y. Preparation and ferroelectric properties of SrBi2Ta2O9 thin films. Appl Phys Lett. 1995;66:221–3. Anpo M, Takeuti M. The design and development of high reactive titanium oxide photocatalysts operating under visible light irradiation. J Catal. 2003;216:505–16. Aravind PR, Shajesh P, Soraru GD, Warrier KG. Ambient pressure drying: a successful approach for the preparation of silica and silica based mixed oxide aerogels. J Sol–gel Sci Technol. 2010;54:105–17. Asahi R, Morikawa T, Ohwaki T, Aoki K, Taga Y. Visible light photocatalysts in nitrogen-doped titanium oxides. Science. 2001;293:269–71. Asakuma N, Fukui T, Toki M, Imai H. Low-temperature synthesis of ITO thin films using an ultraviolet laser for conductive coating on organic polymer substrates. J Sol-Gel Sci Technol. 2003;27:91–5. Avnir D, Levy D, Reisfeld R. The nature of the silica cage as reflected by spectral changes and enhanced photostability of trapped rhodamine 6G. J Phys Chem. 1984;88:5956–9. Azevedo RCS, Sousa RG, Macedo WAA, Sousa EMB. Combining mesoporous silica-magnetite and thermally-sensitive polymers for applications in hyperthermia. J Sol–gel SciTechnol. 2014;72:208–18.

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Tanaka T, Takahashi Y, Iizaka H, Ohguma K, Yamahiro M. Development of super durable hard coat film. In: Proceedings of RadTech Asia, Yokohama. 2011. p. 342–343. Tani T, Namikawa H, Arai K, Makishima A. Photochemical hole-burning study of 1, 4-dihydroxyanthraquinone doped on amorphous silica prepared by alcoholic method. J Appl Phys. 1985;58:3559–62. Three M (3M) Company. Private communication. 1983; Catalogue Nextel fiber. Tillotson TM, Hrubush LW. Transparent ultralow-density silica aerogels prepared by a two-step sol–gel process. J Non-Cryst Solids. 1992;145:44–50. Tohge N, Matsuda A, Minami T. Preparation of ZrO2 and ZrO2-SiO2 coating thin film by sol–gel method. J Chem Soc Jpn. 1987;9:1952–7. Tohge N, Matsuda A, Minami T, Matsuno Y, Katayama S, Ikeda Y. Fine-patterning on glass substrates by the sol–gel method. J Non-Cryst Solids. 1988;100:501–5. Tohge N, Tatsumisago M, Minami T. Preparation of high-Tc superconducting oxide films in the Bi-(Pb)-Ca-Sr-Cu-O system from stabilized metal alkoxides. J Non-Cryst Solids. 1990;121:443–7. Toki M, Miyashita S, Takeuchi T, Kanbe S, Kochi A. A large-size glass produced by a new sol–gel process. J Non-Cryst Solids. 1988;100:479–82. Turniansky A, Avnir D, Bronshtein A, Aharonson N, Altstein M. Sol–gel entrapment of monoclonal anti-atrazine antibodies. J Sol–gel Sci Technol. 1996;7:135–43. Uhlmann DR, Teowee G, Boulton JM, Motekef S, Lee SC. Electrical and dielectric properties of chemically derived ferroelectric films. J Non-Cryst Solids. 1992;147/148:409–23. Uhlmann DR, Dawley JT, Poisl WH, Zelinski BJJ. Ferroelectric films. J Sol–gel Sci Technol. 2000;19:53–64. Vadivel MA, Muraligan T, Manthium A. One-pot microwave-hydrothermal synthesis and characterization of carbon-carbon LiMPO4 (M = Mn, Fe, and Co) cathode. J Electrochem Soc. 2009;156:A79–83. Van den Ham EJ, Peys N, De Dobbelaere C, D’Haen J, Mattelaer F, Detavernier C, Notten PHL, Hardy A, Van Bael MK. Amorphous and perovskite Li3xLa(2/3)- xTiO3 (thin) films via chemical solution deposition: solid electrolytes for all-solid-state Li-ion batteries. J Sol–gel Sci Technol. 2015;73:536–43. Vasanti R, Kaplana D, Renganathan NG. Olivine-type nanoparticle for hybrid supercapacitor. J Solid State Electrochem. 2008;12:961969. Wallace S, Hench LL. The processing and characterization of DCCA modified gel-derived silica. Mater Res Soc Symp Proc. 1984;32:47–52. Wang J, Huang S. Potential of low-temperature post processing of silica for high-temperature stable LED encapsulant. J Sol–gel Sci Technol. 2012;64:557–63. Wu F, Chen KC, Mackenzie JD. Ferroelectric Ceramics – The sol–gel method versus conventional processing. Mat Res Soc Symp Proc. 1984;32:169–74. Wu W, Liang S, Ping Z, Zheng H, Wu L. Low temperature synthesis of nanosized ZnNb2O6 photocatalysts. J Sol–gel Sci Technol. 2012;61:570–6. Xu Y, JihChen C, Xu R, Mackenzie JD. The self-biased heterojunction effect of ferroelectric thin film on silicon substrate. J Appl Phys. 1990;67:2985–91. Yajima S, Okamura K, Hayashi J, Omori M. Synthesis of continuous SiC fibers with high tensile strength. J Am Ceram Soc. 1976;59:324. Yamane M, Aso S, Okano S, Sakaino T. Low temperature synthesis of a monolithic silica glass by the pyrolysis of a silica gel. J Mat Sci. 1979;14:607–11. Yanagisawa T, Shimizu T, Kuroda K, Kato C. The preparation of alkyltrimethyl- ammmoniumKanemite complexes and their conversion to microporous materials. Bull Chem Soc Jpn. 1990a;63:988–92. Yanagisawa T, Shimizu T, Kuroda K, Kato C. The trimethylsilyl derivatives of alkyl- trimethylammmonium-Kanemite complexes and their conversion to microporous SiO2 materials. Bull Chem Soc Jpn. 1990b;63:1535–7. Yokogawa H, Yokoyama M. Hydrophobic silica aerogels. J Non-Cryst Solids. 1995;186:23–9.

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Yokogawa Y, Nishizawa K, Nagata F, Kameyama T. Bioactive properties of chitin/chitosan-calcium phosphate composite materials. J Sol–gel Sci Technol. 2001;21:105–13. Yoldas BE. Alumina sol preparation from alkoxides. Am Ceram Soc Bull. 1975b;54:289–90. Zarzycky J, Prassas M, Phalippou J. Synthesis of glasses from gels: the problems of monolithic gels. J Mater Sci. 1982;17:3371–9. Zhao L, Zhao X, Liu J, Zhang A, Wang D, Wei B. Fabrications of Nb-doped TiO2 (TNO) transparent conductive oxide polycrystalline films on glass substrates by sol–gel method. J Sol–gel Sci Technol. 2010;53:475–9. Zhu CM, Wang LG, Yuan SL, Tian SC. Room-temperature multiferroic properties of 0.6BiFeO30.4(Bi0.5Na0.5)(1-x)BaxTiO3 solid solution ceramics. J Sol–gel Sci Tech. 2015;76:289–97. Zhuang H, Kozuka H, Yoko T, Sakka S. Preparation of superconductive Bi-Sr-Ca- Cu-O coating films by the sol–gel method using an aqueous solution of metal acetate. Jpn J Appl Phys. 1990;29:L1107–10. Zieba J, Zhang Y, Prasad PN. Sol–gel processed inorganic oxides organic polymer composites for second-order nonlinear optical application. SPIE Sol–gel Optics II. 1992;1328:403–9.

2

The Synthesis and Solution Stability of Alkoxide Precursors Vadim G. Kessler

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Synthesis of Homometallic Precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Reactions of Metals with Alcohols (Method 1.1) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Anodic Oxidation of Metals (Method 1.2) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Reactions of Metal Oxides or Hydroxides with Alcohols (Method 1.3) . . . . . . . . . . . . . . . . . . . . Alcoholysis of Metal Derivatives of Weak or Volatile Acids (Method 1.4) . . . . . . . . . . . . . . . . . Metathesis Reactions with Alkali Alkoxides and Ammonia or Amines (Method 1.5) . . . . . Alcohol Interchange Reactions (Method 1.6) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Redox Processes in Approach to Alkoxide Precursors (Method 1.7) . . . . . . . . . . . . . . . . . . . . . . . Synthesis of Heterometallic Alkoxide Precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Heterometallic Alkoxides Formed via Lewis Acid-Base Interaction . . . . . . . . . . . . . . . . . . . . . . . . Heterometallic Alkoxides Formed via Formation of Heteronuclear Metal-Metal Bonds or Isomorphous Substitution . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Solution Stability with Respect to Formation of Oxoalkoxides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Partial Hydrolysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Oxidation by Oxygen from Atmosphere or Dissolved in Solvents . . . . . . . . . . . . . . . . . . . . . . . . . . Ether and Ester Elimination . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . β-Hydrogen Elimination . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermal Desolvation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Solution Stability with Respect to Solvolysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Individual Alkoxide Complexes Applied as Precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Homometallic Precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Heterometallic Precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . The Chemistry of Alkoxide Precursors’ Transformation into Oxide Materials in Sol-Gel Technology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

32 33 33 35 36 37 39 40 42 42 43 47 48 48 49 50 51 52 52 54 54 55 64 69

V. G. Kessler (*) Department of Chemistry and Biotechnology, Swedish University of Agricultural Sciences (SLU), Uppsala, Sweden e-mail: [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_1

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Abstract

The aim of this chapter is to serve as a guide in understanding the principles in the chemical approaches to and stability of the metal and silicon alkoxide precursors. The major accent is made on the overview of the accessible precursor compounds applicable in sol‐gel technology, focusing on their synthesis and reactivity including the chemistry behind the sol‐gel process. The chapter provides an insight into the synthesis of homometallic precursors (see section “Synthesis of Homometallic Precursors”), synthesis of heterometallic precursors (see section “Synthesis of Heterometallic Alkoxide Precursors”), solution stability with respect to formation of oxoalkoxides (see section “Solution Stability with Respect to Formation of Oxoalkoxides”), solution stability with respect to solvolysis (see section “Solution Stability with Respect to Solvolysis”), and, finally, a short review summarizing the literature data on individual alkoxide complexes applied as precursors (see section “Individual Alkoxide Complexes Applied as Precursors”) considering separately the homometallic species (see section “Homometallic Precursors”) and the heterometallic ones – so-called singlesource precursors (see section “Heterometallic Precursors”). Finally, a brief overview of the modern concepts treating the transformation of precursors into materials is also provided (see section “The Chemistry of Alkoxide Precursors’ Transformation into Oxide Materials in Sol-Gel Technology”).

Introduction The development of sol‐gel technology has at a very early step put forward a request on development of precursor compounds – chemical substances that have high solubility in organic solvents are easily transformed into chemically reactive forms of hydrated oxides on hydrolysis. They should display considerable stability in solution to guarantee the reproducibility of the materials preparation and, last but not the least, be easy to be purified to provide sufficient chemical quality of the final products. Metal alkoxides, M(OR)n, are derivatives of alcohols, ROH, which are usually easily accessible and inexpensive organic compounds and are extremely weak as acids, easily removable via hydrolysis and thermal treatment, leaving high purity hydrated oxides. This circumstance made metal alkoxides the most common candidates for the role of molecular precursors (Veith 2002; Jones 2002; HubertPfalzgraf 2003; Kessler 2003). The works in this field during the last 20 years, including both the studies of the molecular and crystal structure and the reactivity of these compounds, have considerably changed their image in the eyes of both chemists and the materials scientists: it turned out that sometimes the compounds that are the most stable products in the reactions of synthesis of metal alkoxides and that were earlier considered to be M(OR)n are in fact oxoalkoxides MOx(OR)y. In many cases, especially for the preparation of complex solutions, including derivatives of several metals, it turned out impossible to use only the derivatives of aliphatic alcohols, CnH2n+1OH, because of their poor solubility, stability, or reactivity. This gave

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33

rise to the development of two new types of alkoxide precursors – derivatives of functional alcohols (alkoxyalcohols and aminoalcohols), on one hand, and heteroleptic alkoxides, including other ligands (such as β-diketonate, carboxylate, and aminoalkoxide ones) in addition to common aliphatic alkoxide groups, on the other. This change in the understanding of nature of alkoxide complexes has been reflected even by the titles of the modern textbooks on this topic called “Alkoxide and Phenoxide Derivatives of Metals” (Bradley 1965; Bradley et al. 2001) and “The Chemistry of Metal Alkoxides” (Turova et al. 2002). The complexity of situation has been increased even more by the rise of a still quite small but quickly growing family of alcoholates – highly soluble complexes of metal carboxylates or β-diketonates with functional alcohols. The latter do not contain formally the alkoxide ligands but are related to metal alkoxides in many of their properties and find the increasing application in sol‐gel technology. The aim of this chapter is to serve as a guide in understanding the principles in the chemical approaches to and stability of the metal alkoxide precursors known already in fact for all the elements of the periodic table, excluding only the highly radioactive ones. The major accent is made on the laboratory approaches to the soluble and chemically reactive alkoxide derivatives, applicable in sol‐gel technology. The chapter includes synthesis of homometallic precursors (see section “Synthesis of Homometallic Precursors”), synthesis of heterometallic precursors (see section “Synthesis of Heterometallic Alkoxide Precursors”), solution stability with respect to formation of oxoalkoxides (see section “Solution Stability with Respect to Formation of Oxoalkoxides”), solution stability with respect to solvolysis (see section “Solution Stability with Respect to Solvolysis”), and, finally, a short review summarizing the literature data on individual alkoxide complexes applied as precursors (see section “Individual Alkoxide Complexes Applied as Precursors”) considering separately the homometallic species (see section “Homometallic Precursors”) and the heterometallic ones – so-called single-source precursors (see section “Heterometallic Precursors”). The custom synthesis is a challenging task and requires application of anhydrous solvents and, what is absolutely crucial, an inert moisture-free atmosphere. A proper infrastructure including Schlenk lines of glove boxes needs to be available for success of this work. It has to be noted that a broad variety of precursors are commercially available nowadays (for details, please, see section “Individual Alkoxide Complexes Applied as Precursors”), but need to be stored and handled properly to succeed with their application in sol‐gel preparation of materials.

Synthesis of Homometallic Precursors Reactions of Metals with Alcohols (Method 1.1) Direct Reaction in Inert Atmosphere (Argon or Nitrogen) Direct reaction with alcohols with evolution of hydrogen gas and formation of metal alkoxides is possible only for the most electropositive metals such as alkali, magnesium and alkaline earth, rare earth metals, and aluminum:

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M þ nROH ! MðORÞn þ n=2 H2 ðgÞ

(1)

In fact the reaction proceeds readily enough at room temperature only for the alkali metals and most acidic alcohols such as MeOH and EtOH or the functional ones such as 2-methoxyethanol. Heating is usually indispensable to lead reaction to the completion even in the case of alkaline and alkaline earth metals. The readiness of an alcohol to react with metals may depend strongly on its purity. Thus, comparably small water contents, such as less than 0.5 wt%, may, for example, leave ethanol almost inert in reaction with magnesium or calcium even on reflux, while ethanol with less than 0.01 wt% water has an observable reactivity toward these metals even at room temperature. To increase the reaction temperature in the case of magnesium or rare earth metals, the reaction is often carried out in a mixture of a parent alcohol (usually iPrOH in the latter case) and a high-boiling point hydrocarbon (toluene or even xylene, in 1:1 or 1:2 volume ratio to alcohol). The reaction requires the use of a catalyst for the alkaline earth metals, rare earth metals, and aluminum. The most common approaches are the use of (in the laboratory practice only) the salts of mercury(II) such as HgCl2 or Hg(OAc)2. Very small portions of these salts cause amalgamation of the metal surface (and thus clean it from the oxide layer) and facilitate the reaction with alcohols. The larger-scale synthesis (and thus the industrial one – in the scope of pollution danger) uses the initial addition of solid iodine (1 g or less per 100 g of alkoxide to be prepared). Formation of metal iodide serves both for cleaning the surface and increases also slightly the acidity of alcohols via formation of solvate complexes. In the case of barium, the application of dry ammonia gas has been reported for this purpose (Caulton et al. 1990; Drake et al. 1992). The major factor facilitating the reaction of metals with alcohols is the solubility of the alkoxides formed. Insoluble alkoxides form a protective layer on the surface of the metal and it hinders the reaction. Even the reaction of sodium with tBuOH in toluene may be almost stopped by the formation of poorly soluble NaOtBu. It is to be mentioned that the reaction of metals with an excess of alcohol leads normally, except for aluminum, to formation of solvates with corresponding alcohols, such as Li(OEt)2EtOH, Mg(OMe)23.5MeOH, Ca(OEt)24EtOH, [Nd(OiPr)3 (iPrOH)]4, etc. (see 1.5). To obtain non-solvated species, the reaction should be carried out with a stoichiometric amount of alcohol in a different solvent (most often, toluene) (Fisher and McElvain 1934). It is also important to notice that the reaction products quite often contain impurities of oxoalkoxides resulting from the presence of residual oxygen in the solvents or quite complicated redox side reactions. If such oxoalkoxides possess considerable thermodynamic stability, as, for example, Ba6O (OC2H4OMe)10 (MeOC2H4OH)4 (Caulton et al. 1993) or Ln5O(OiPr)13, where Ln = Sc, Y, or lanthanides (Hubert-Pfalzgraf et al. 1997a), their formation cannot be avoided, and they will in any case be isolated as the major reaction product and may be purified further by recrystallization. The reaction of metals with alcohols in inert atmosphere (except for the alkali ones) leaves very often a dark residue of unreacted small particles of metal or metal suboxides. This kind of residue is almost inevitable for aluminum and rare earth metals and can be simply removed by decantation at the end of the reaction.

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35

Oxidation of Metals by Oxygen Gas in Alcohol Media This approach offers only extremely highly soluble and stable alkoxide complexes with rather high resistance to hydrolysis. It was first applied to the preparation of thallium(I) ethoxide, carried out in a Soxhlet filter: 2Tl þ EtOH þ 1=2 O2 ! TlOEt þ TlOH

(2)

The reaction results in formation of a double-layer system, where the bottom layer is a 95 wt% solution of TlOEt in EtOH, while the upper one is containing almost all TlOH (Turevskaya et al. 1975). Oxidation by oxygen in air and solvents turned also to be a useful tool in approach to the copper(II) derivatives of aminoalcohols. The hydrolytic stability of Cu (II) aminoalkoxides appears to be its driving force. It is also apparently the same for the reaction of copper metal with alcohols in the presence of N-donor ligands, L, and considerable excess of alcohol (Kovbasyuk et al. 1998): Cu þ excess ROH þ nL þ 1=2 O2 ! CuðORÞ2  nL  mROH þ H2 O

(3)

Anodic Oxidation of Metals (Method 1.2) Considered initially as, in general, a simple extension of the direct reaction with alcohols for less active metals by application of an anodic potential, the anodic oxidation of metals turned to be a much more complicated process. At present, at least three different oxidation mechanisms have been proposed for different groups of metals: – The most active metals, such as lanthanides, receive really just a support for the direct interaction with alcohols (2-propanol in this case) from the applied anodic potential supposedly via the elimination of the oxide barrier. The electric current yields (the ratio of the alkoxide obtained to the total charge that passed through the system) often exceed 100%. High concentrations of soluble conductive additives (LiX or R4NX, where X = Cl, Br), which contaminate the product have to be removed by repeated recrystallization from hydrocarbon solvents. – The late transition and main group metals follow the anodic oxidation pathway analogous to that in aqueous solutions. The minimal oxidation potentials in these cases can in fact be very low (up to max. 3.0 V), while the higher ones are readily applied to accelerate the process. The anodic reaction consists of dissolution of metal ions in the form of anionic halide complexes, which are later transformed into insoluble alkoxides by reaction with alkoxide anions generated at the cathode, for example (Lehmkuhl et al. 1975): Cathode : ROH þ e ! RO þ 1=2 H2 ðgÞ Anode : Cu þ 4Cl  2e ! CuCl4 2



(4) (5)

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Solution : CuCl4 2 þ 2RO ! CuðORÞ2 ðsÞ þ 4Cl

(6)

Only insoluble alkoxides can be obtained by this method because the soluble ones are normally reduced at the cathode, transforming the process into the electrochemical transport of the metal from anode to cathode. The products again are usually heavily polluted by halide admixtures and should be then washed repeatedly with alcohols to remove adsorbed conductive additives (Hubert-Pfalzgraf et al. 1997b). It has, however, been reported that application of amines (such as dipyridyl, phenantroline), giving rather stable insoluble complexes with Cd and Cu alkoxides, allows alkoxides free from halide admixtures to be isolated (Banait and Pahil 1986). – The early transition metals are dissolved via a complex mechanism involving oxidation of alkoxide ligands with formation of extremely reactive alkoxoradicals that in turn attack the metal, forming soluble alkoxide complexes already at the anode: M þ n RO  n e ! MðORÞn

(7)

The reaction has highest speed in the alcohols displaying highest electric conductivity, such as MeOH or MeOC2H4OH. Low concentrations of conductive additives applied in this case assure high purity of the final product. It is in fact very important to keep the concentrations of the additives in the interval 0.01–0.05 M as the high potentials applied cause the formation of free halogens that oxidize the alcohols and provide finally water as by-product, leading to the formation of oxoalkoxide impurities. The other impurity formed simultaneously is the alkoxide derived from the conductive additive, for example, lithium alkoxides from lithium halides. On interaction with the metal alkoxide, they provide heterometallic complexes. Thus, a whole series of different bimetallic Li-Mo and Li-W alkoxides have been isolated and characterized as by-products of the electrochemical syntheses of M(VI) alkoxides (Kessler et al. 1998b). Another source of oxoalkoxide impurities is the cathodic reduction, which transforms low oxidation state impurities into oxoalkoxides via subsequent re-oxidation by oxygen dissolved in solvents. Following the optimized procedures, it is possible, however, to produce rather high-quality methoxide derivatives of Nb (Turevskaya et al. 1995a, b, c), Ta (Turova et al. 1996), Mo (Kessler et al. 1993), W (Seisenbaeva et al. 2001a), and Re (Seisenbaeva et al. 2001b).

Reactions of Metal Oxides or Hydroxides with Alcohols (Method 1.3) This reaction is useful for preparation of alkoxides from most basic or most acidic oxides and hydroxides. The alkoxides obtained should have quite high hydrolytic and thermal stability, because water formed during the reaction is removed by distillation as an azeotrope with an aromatic hydrocarbon solvent (usually toluene).

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In the laboratory practice, it can be applied for the preparation of phenoxides of alkali or alkaline earth metals, for example: NaOH þ PhOH ! NaOPh þ H2 O

(8)

Recent crystal structure studies have shown that the interaction of basic hydroxides with aliphatic alcohols does not lead to metal alkoxides but to alcohol solvates of the hydroxides. For example, the reaction of Ba(OH)2 with MeOH was found to provide Ba(OH)2(MeOH)2MeOH (Turova priv. comm.). This reaction has been in contrast successfully applied for the synthesis of alkoxide derivatives of acidic oxides, as the whole homologous series of vanadium alkoxides (Orlov and Voronkov 1959; Prandtl and Hess 1913), and for the preparation of a number of hydrocarbon-soluble complexes with diols of molybdenum (VI) (Bishop et al. 1979), rhenium(VI) (Edwards et al. 1998), and osmium (VI) (Lehtonen et al. 1999): V2 O5 þ 6 ROH ! 2VOðORÞ3 þ 3H2 O, R ¼ Et  C5 H11 MoO3 þ 3½MeCHðOHÞ2 ! ðMeCHOÞ2 MoO2  2RðOHÞ2 þ H2 O

(9) (10)

Re2 O7 þ 4HOCMe2 CMe2 OH ! 2ReO3 ðOCMe2 CMe2 OHÞðHOCMe2 CMe2 OHÞ (11) The reaction achieves completeness, when the alkoxides thus formed display quite considerable resistance to hydrolysis and can then be purified by some efficient technique (distillation for vanadium derivatives and recrystallization for the diolates).

Alcoholysis of Metal Derivatives of Weak or Volatile Acids (Method 1.4) The action of alcohols on the metal derivatives of extremely weak and, which is of special importance, highly volatile acids, for example, alkyls, carbides, nitrides, amides, alkyl amides, silazides, hydrosulfides, hydrides, etc., provides an approach to high purity samples of metal alkoxides, usually under extremely mild conditions. The reaction MXn þ nROH ! MðORÞn þ nHX X ¼ H, Alk, C= 2C, 2 = N, NH2 , NR2 , NðSiR3 Þ2 , SH

(12)

is usually carried out in a volatile hydrocarbon solvent (such as hexanes or pentane) and the products are purified by evaporation of the byproducts and the solvent in vacuum, leaving the target alkoxide as the residue. Hydrides can be used as sources of alkaline metal alkoxides (LiH, NaH) in the reactions with halogenated alcohols, such as, (CF3)3COH, to avoid the danger of condensation of Wurtz type (Dear

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et al. 1970). Metal alkyls have been applied earlier for the preparation of a number of early transition metal derivatives, for example, CuOMe (Costa et al. 1965), Cr(OR)2 (Chisholm et al. 1979), and V(OtBu)4 (Razuvaev and Drobotenko 1977), but are themselves extremely unstable and normally not commercially available, which precludes their application in common laboratory practice. The most broadly applied laboratory approach of this type is the reaction of metal alkyl amides, usually bis-(trimethylsilyl)-amides with the stoichiometric amounts of alcohols. The starting reagents even in this case are not available commercially but can be obtained more or less easily by reaction of the corresponding metal chlorides with commercially available LiN(SiR3)2 in anhydrous diethyl ether:
 MCln þ nLiNðSiR3 Þ2 ! M NðSiR3 Þ2 n þ nLiCl

(13)

M[N(SiR3)2]n can then be purified – for the main group derivatives (for application in the synthesis of alkoxides, see Zn (Goel et al. 1990a), Cd (Boulmaaz et al. 1992), Pb (Matchett et al. 1990; Papiernik et al. 1989), and Bi (Massiani et al. 1990; Goel et al. 1990a)) – by sublimation direct from the reaction mixture, after removal of the Et2O in vacuum, and (for the early transition metal compounds (Cr(II), Mn (II) (Horvath et al. 1979)), after the removal of ether) by the extraction from the residue with pentane or hexanes, separating LiCl by decantation. It should be mentioned that this approach is hardly practically applicable for the synthesis of the derivatives of late transition metals such as Co, Ni, or Cu because of poor stability of their amide derivatives (Bryndza and Tam 1988). It should be mentioned that the reaction of metal chlorides with alcohols could not be applied for the synthesis of metal alkoxides – precursors of oxide materials. Its products are usually quite complex mixtures of alkoxide chlorides and alcohol solvates of metal oxochlorides (Turova et al. 2002; Turevskaya et al. 1989). Formation of alcoholates – solvate complexes with functional alcohols – can be considered as a variety of this synthetic approach. Metal β-diketonates or carboxylates are reacted with amino- or alkoxy-alcohols in stoichiometric amounts in organic solvents (both nonpolar, such as toluene or hexane, or polar, such as methanol or ethanol, can be applied (Williams et al. 2001; Seisenbaeva to be published): NiðOAcÞ2  4H2 O þ 2HOC2 H4 NMe2 ! NiðOAcÞ2 ðHOC2 H4 NMe2 Þ2 þ 4H2 O

(14)

MnðacacÞ2  xH2 O þ 2HOCHMeCH2 NMe2 ! MnðacacÞ2 ðHOCHMeCH2 NMe2 Þ

(15)

The advantage of the alcoholate complexes lies in their high solubility in organic solvents. They provide also a possibility to avoid more complicated dehydration procedures necessary for the derivatives of late transition metals to be used for the preparation of complex solutions together with metal alkoxides.

2

The Synthesis and Solution Stability of Alkoxide Precursors

39

Metathesis Reactions with Alkali Alkoxides and Ammonia or Amines (Method 1.5) This approach remains the most commonly applied in the synthesis of metal alkoxides. The starting reagents are the anhydrous metal halides, most often chlorides, or other anhydrous metal salts, such as nitrates or acetates: MXn þ nROH þ nR3 N ! MðORÞn þ nR3 NHX

(16)

MXn þ nMI OR ! MðORÞn þ nMI X

(17)

The traditional technique using ammonia gas has been applied for the preparation of the alkoxides of titanium (Demarcay 1875), zirconium and hafnium (Bradley et al. 1952), cerium (IV) (Bradley and Holloway 1962), and niobium and tantalum (Bradley et al. 1956a, b). In this approach, a halide or a pyridinium halogenometallate salt, for example, (PyH)2ZrCl6, is dissolved in a mixture of toluene with the parent alcohol, and ammonia gas is bubbled through the solution for several hours. The voluminous precipitate of NH4Cl is removed by filtration and washed with the alcohol on the filter to improve the yield of the soluble alkoxide. To avoid the use of ammonia gas and simplify the procedure as a whole and specifically the separation of the ammonium salts, there has been proposed to use the amines such as triethylamine or pyridine in the same purpose. This route provided access, for example, to the stable samples of MoO(OiPr)4 from MoOCl4 (Chisholm et al. 1984) and those of Re2O3(OMe)6 from ReOCl4 (Edwards et al. 1980). It is necessary to mention that neither ammonia nor amines can be applied for the preparation of pronouncedly basic alkoxides – derivatives of alkali, alkaline earth, or rare earth metals (their formation is impossible in the presence of acidic ammonium salts). A specific problem in application of ammonia or amines lies in the need of introducing a metal halide into this reaction as a solution in a solvent mixture including the parent alcohol. Strong Lewis acids such as metal halides are at room temperature prone to convert the alcohols, especially the ramified ones, into alkyl halides and transforming themselves into oxohalides (Turova et al. 2002). This side reaction decreases the yield of the target products and, when the tertiary (Bradley et al. 1978) or aromatic (Niederberger et al. 2002) alcohols are used, can lead (in not completely anhydrous conditions) even to formation of oxides or hydrated oxides. This effect can be avoided if the halides are introduced as solutions in aprotic solvents (toluene, ether, THF) into the solutions of alkali alkoxides, for example: FeCl3 =toluene þ 3NaOEt=EtOH ! FeðOEtÞ3 þ 3NaCl

(18)

BiCl3 =ether þ 3NaOEt=EtOH ! BiðOEtÞ3 þ 3NaCl

(19)

Strong cooling is always recommended at the initial step of this process. Then the reaction mixtures are usually warmed to room temperature after the complete

40

V. G. Kessler

addition of the halide and then even often subjected to reflux in order to destroy the possible heterometallic impurities. The heterometallic impurities are sometimes so stable, for example, NaZr2(OR)9, that they can even be distilled in vacuum without decomposition (Bartley and Wardlaw 1958). It is important to avoid the possible deviations in the reaction stoichiometry, but even the perfect one does not guarantee the purity of the obtained samples because the alkoxide chlorides or bimetallic alkoxide chlorides can sometimes display really high stability. For example, Y3(OtBu)8Cl2THF or Nd6(OiPr)17Cl have been isolated as the major products in the reaction of the corresponding trichlorides with three equivalents of NaOR (Evans et al. 1988; Andersen et al. 1978). In many cases, larger halide ligands (Br or I instead of Cl) or a larger alkaline metal atoms (K instead of Na or Li) can help to avoid the side reactions of this kind, for example, in Kritikos et al. (2001),   5LnCl3 þ 15KOi Pr þ H2 O ! Ln5 O Oi Pr 13 þ 2iPrOH

(20)

Metal acetates have been used as reagents in the metathesis with alkali alkoxide mainly in order to produce the main group metal derivatives, for example, the homologous series of lead alkoxides (Papiernik et al. 1989; Turevskaya et al. 1982). The reaction produces the insoluble sodium acetate that can be removed by filtration or decantation. It should be mentioned that, when carried out in toluene (on reflux), it could very easily lead to oxoalkoxide derivatives via ester or ether elimination side reactions (see section “Synthesis of Heterometallic Alkoxide Precursors”). Application of the nitrate complexes has been proposed in the metathesis-based approaches to the derivatives of Ce(IV) in the view of their much higher stability and commercial availability (Gradeff et al. 1985). It is necessary to mention that during the development of metathetic approaches, a number of alkoxylating agents other than the alkali alkoxides have been tested in this purpose. For example, the gas phase co-condensation of volatile metal fluorides or chlorides with alkylsiliconalkoxides has been reported for the preparation of M (OMe)6, M = Mo, W, and Re (Jacob 1982; Bryan et al. 1991). This technique requires a special equipment and provides rather small quantities of these products that can be obtained much more easily by the anodic oxidation of the corresponding metals. Another example of a different alkoxylating agent is the soluble Mg(OMe)2, which has been used to produce the methoxides from the metal fluorides (Bryan et al. 1991).

Alcohol Interchange Reactions (Method 1.6) The equilibrium reactions of metal alkoxides with alcohols: MðORÞn þ R’OH Ð MðORÞn1 ðOR’Þ þ ROH;

(21)

2

The Synthesis and Solution Stability of Alkoxide Precursors

41

are often very flexible. It is important to keep this fact in mind in developing the procedures for sol‐gel applications of the alkoxide precursors: when dissolved in other alcohols than the parent one, they would undergo a ligand exchange that can change their molecular structure and hydrolytic properties and, in case of the heterometallic compounds (to be used as single-source precursors), even their chemical compositions. This method is useful as a synthetic procedure, when an efficient synthetic approach or commercial availability provides a different homologue than the one desired for further applications. The completeness of the reaction is achieved more easily if the desired product possesses much lower solubility in the new parent alcohol, for example:     ZrðOn PrÞ4 þ excessi PrOH ! ZrðOn PrÞ Oi Pr 3 i PrOH

(22)

or if the alcohol to be introduced has a considerably higher boiling point, which facilitates the removal of the other alcohol, formed in the reaction, by vacuum distillation, for example (Johansson et al. 2000):   MoOðOMeÞ4 þ excessi PrOH ! MoO Oi Pr 4 þ 4MeOH

(23)

The treatment with the new alcohol must be repeated several times (with complete dissolution of the crude intermediate product) to insure the completeness of transformation. It should be mentioned that the completeness in the exchange of the alkoxide groups might not always be achieved: the stability of molecular structures including small bridging ligands or functional chelating alkoxide ligands prevent in many cases the possibility to replace them, for example (Bradley et al. 1978; Johansson et al. 2000):   TaðOMeÞ5 þ excessi PrOH ! TaðOMeÞ Oi Pr 4 þ 4MeOH MoO2 ðOC2 H4 OMeÞ2 þ excess EtOH ! recrystallization without substitution (24) Replacement can also be achieved by other sources than alcohols, for example, esters. This is of interest if the boiling points of the two alcohols are very close (e.g., i PrOH and tBuOH, b.p. 82.4  C) or when the alcohol to be applied is highly unstable (e.g., silanols, unsaturated alcohols, etc.):   Zr Oi Pr 4 þ 4AcOt Bu ! ZrðOt BuÞ4 þ 4AcOi Pr

(25)

  Ti Oi Pr 4 þ 4AcOSiR3 ! TiðOSiR3 Þ4 þ 4AcOi Pr

(26)

The reactions with esters are most often carried out in aromatic hydrocarbon solvents to decrease the reaction temperature by removing a more volatile azeotropic mixture of a new ester with, for example, toluene or (in the past) benzene (Bradley et al. 1978; Bradley and Thomas 1959).

42

V. G. Kessler

A specific class of the ligand exchange reactions and, in some cases, of alcohol interchange improves the solubility and behavior in the hydrolysis and subsequent gelation processes. The reaction can in general be written as follows: MðORÞn þ mHZ ! MðORÞnm Zm þ mROH

(27)

where HZ represents aminoalcohols or other functional alcohols, β-diketones, or carboxylic acids. These reactions have been rather thoroughly studied for the derivatives of M(IV) such as titanium, zirconium, and cerium and are described in a number of recent review articles (Ribot et al. 1991; Hubert-Pfalzgraf 2003; Jones 2002).

Redox Processes in Approach to Alkoxide Precursors (Method 1.7) The redox reactions do not in fact belong to the common approaches in preparation of precursors for the sol‐gel technology. The only example worth noting here is the oxidation of low-valent chromium derivatives (dibenzene-chromium, Cr(OR)3) by the t-butylperoxide, providing access to chromium(IV) alkoxides – highly soluble and volatile compounds (Krauss and Munster 1967). On the other hand, the redox reactions are in many cases responsible for the transformation of metal alkoxides in solutions leading to formation of oxoalkoxides and will be discussed below in section “Solution Stability with Respect to Formation of Oxoalkoxides.”

Synthesis of Heterometallic Alkoxide Precursors The special interest in heterometallic alkoxide complexes is due to the possibility of their application as single-source precursors in the preparation of complex inorganic materials (oxides, sulfides, metal alloys, and even nanocomposites) (Veith 2002). The single-source precursor represents a compound containing the necessary elements in desired stoichiometric ratio. Synthesis and properties of the heterometallic alkoxides have been described in detail in a number of recent reviews (Veith 2002; Jones 2002; Hubert-Pfalzgraf 2003; Caulton et al. 1990; Kessler 2003). The formation of heterometallic complexes in general can occur due to one of the three following factors: 1. Lewis acid-base interaction (exploiting the difference between two or several metal atoms in electronegativity, which permits to consider one metal center as a stronger acceptor of the electron density and the alkoxide or other ligands at the other as a better donor of it) 2. Formation of a heterometallic metal-metal bond, which in this case should also provide a donor-acceptor interaction

2

The Synthesis and Solution Stability of Alkoxide Precursors

43

3. Isomorphous substitution, which might in some cases not lead to formation of the true heterometallic species, but provides in any case the homogenization at the molecular level The synthetic approaches to heterometallic complexes will be classified here below according to these three principles providing their formation.

Heterometallic Alkoxides Formed via Lewis Acid-Base Interaction Complex Formation Between Two Alkoxides (Method 2.1) The pronounced Lewis basicity of the alkoxide ligands of the alkali and alkaline earth metal alkoxides explains their capacity to form heterometallic complexes in solution with the vast majority of high-valent transition or main group metal alkoxides, for example: LiOR þ NbðORÞ5 ! LiNbðORÞ6

(28)

The chemical composition of the products is determined by a number of important factors, such as the nature of metal atoms involved, the nature of alkoxide groups, the ratio of homometallic reactants applied, and, in certain cases, even on the solvent in which the reaction is carried out. For example, the reaction of barium ethoxide and titanium ethoxide in 1: x (where x  2) ratio in oxygen-free solutions provides different products in alcohol and in hydrocarbon media (Yanovsky et al. 1995; Kessler et al. 1994a): EtOH

BaðOEtÞ2 þ > 2TiðOEtÞ4 ! BaTi2 ðOEtÞ10 ðEtOHÞ5 toluene

BaðOEtÞ2 þ > 2TiðOEtÞ4 ! BaTi4 ðOEtÞ18

(29) (30)

Formation of different complexes at different ratios of reactants in the same solvent (THF in this case) can be illustrated by the examples from barium-zirconium isopropoxide system (Vaartstra et al. 1991) (the authors used barium metal in the presence of solvating alcohol, but application of the ready alkoxide gives the same result (Turevskaya et al. 1995a)): h       i Ba þ Zr2 Oi Pr 8 i PrOH 2 ! BaZr2 Oi Pr 10 þ H2 ðgÞ 2

h       i Ba þ 2Zr2 Oi Pr 8 i PrOH 2 ! Ba Zr2 Oi Pr 9 þ H2 ðgÞ þ 2i PrOH 2

(31) (32)

Synthesis of heterometallic complexes by direct interaction of homometallic alkoxides has been reported in many cases also for the rare earth metals, but in this case, it is necessary to keep in mind that the commercial “Ln(OiPr)3” usually

44

V. G. Kessler

contain the oxoalkoxide complex, Ln5O(OiPr)13, as their major component. The reactivity of the latter toward other alkoxides is comparably low, and prolonged refluxing in toluene or the reaction in a melt is recommended to insure the completeness of transformation (Poncelet et al. 1989a, b). The only reaction between the two high-valent metal alkoxides, not involving specific mechanisms with formation of oxoalkoxides, is the formation of the aluminum and hafnium isopropoxide (Turevskaya et al. 1997):         Hf Oi Pr 4 i PrOH þ 2Al Oi Pr 3 ! HfAl2 Oi Pr 8 þ 2i PrOH

(33)

It is also important to notice that this reaction takes place even in the alcohol media, but gives in this case far not quantitative yields of the product, whose formation took several days.

Metathesis of a Metal Halide with a Bimetallic Alkoxide of Another Metal and an Alkali Metal (Method 2.2) This approach has been proposed for the case, where the alkoxide of one of the metals is not easily accessible or is an insoluble and inert solid as, for example, the alkoxides of late transition and some main group metals (Mn(II), Fe(II), Fe(III), Co, Ni, Cu, Zn, Cd, Sn(II), Pb(II)). It has been applied also for the preparation of heterometallic derivatives of rare earth elements (see Bradley et al. 1978, 2001; Turova et al. 2002). It is necessary to take into account that the reaction:
 MXn þ MI M’ðORÞmþ1 ! M M’ðORÞmþ1 n þ nMI X;

(34)

not always follows the simplified reaction formula given by the Eq. 34. The so-called alkoxometallate ligands, existing in the structures of heterometallic alkoxides of alkali metals, such as [Al(OR)4], [Zr2(OR)9], or [Nb(OR)6], are not the ultimately thermodynamically stable ionic units. They are therefore very rarely just transferred by this reaction from an alkali metal cation to a less electropositive metal cation. The deviations may occur both due to formation of more stable oxocomplexes (Boulmaaz et al. 1994) and even more stable homoleptic alkoxide complexes (Kessler et al. 1994b; Yanovsky et al. 1994), for example:     ZnI2 þ 2NaTa Oi Pr 6 ! Zn2 Ta4 O4 Oi Pr 16       LaCl3 þ 3NaNb Oi Pr 6 ! LaNb2 Oi Pr 13 þ Nb Oi Pr 5 þ 3NaCl

(35) (36)

A serious problem that can also be associated with application of this technique lies in the possibilities of formation of by-products including the halides or alkali metals (or even both as in case of formation of [NaPb2Ti2O(OiPr)10]+Cl (Hubert-Pfalzgraf private comm.). A family of the techniques described below is based on the introduction of new ligands such as oxo-groups or different organic residues such as β-diketonate,

2

The Synthesis and Solution Stability of Alkoxide Precursors

45

carboxylate, or aminoalkoxides groups. Successful development of new approaches of this kind needs prediction of the chemical composition and even the structure of the new complexes to be prepared and choice of the proper reaction stoichiometry. This prediction can be done using the Molecular Structure Design Concept, described in detail by Kessler (2003), and including the following steps: 1. Choice of the structure type to be used (see Fig. 1) 2. Calculation of the necessary number of the donor atoms 3. Choice of the ligands with proper composition and sterical requirements that provide both the right number of donor functions and the protection of the chosen core of metal and donor atoms (placing the metal atoms into the thermodynamically preferred coordination polyhedra) The most complete classification of the stable structure types for the alkoxide complexes can be found in Turova et al. (2002). The practical principles in application of this concept lie either in decreasing the number of donor atoms by replacing two OR groups with an oxo-ligand and thus increasing the strength of the Lewis acids involved or by providing some additional donor atoms from bidentate (usually) chelating ligands, which are necessary to support the chosen structure type (Bradley and Holloway 1965).

Fig. 1 Schematic views of the most important building blocks in the structures of metal alkoxide aggregates

OR

OR M

OR

M

M OR

M

OR

M

M

OR

OR

OR

OR X

M

X

M

OR

OR

M

M

OR

OR

OR

M M

M

OR OR

OR OR

M OR

M

M

OR

M

OR M

OR M

OR

OR

46

V. G. Kessler

Microhydrolysis of Alkoxides of Different Metals in Solutions (Method 2.3) This approach has been first applied to access to the heterometallic alkoxides of bismuth, because the homoleptic Bi(OR)3 usually does not form any heterometallic complexes with the alkoxides of other high-valent metals (Parola et al. 1997): BiðORÞ3 þ 2TiðORÞ4 þ H2 O ! BiTi2 OðORÞ9 , R ¼ Et, i Pr

(37)

This reaction was carried out successfully in an alcohol media. It is necessary to mention, however, that alcohols very often destroy the heterometallic complexes derived from elements with close Lewis acidity. The efficiency of microhydrolysis as an approach to heterometallic species is much better in inert media (hydrocarbon solvent).

Micropyrolysis of Alkoxides of Different Metals in Solutions (Method 2.4) When an alkoxide derivative can easily decompose thermally in solution forming an oxoalkoxide (see section “Solution Stability with Respect to Formation of Oxoalkoxides”), this reaction might be exploited in approach to heterometallic species. This reaction is useful especially for the synthesis of heterometallic derivatives of molybdenum (Johansson et al. 2000; Johansson and Kessler 2000a, b) and zirconium (Kessler et al. 1998a) (see also in 2.5), for example:       4MoO Oi Pr 4 þ 2Ta Oi Pr 5 ! Mo4 Ta2 O8 Oi Pr 14 þ 2ðCH3 Þ2 CO   þ 2i PrOH þ 4 i Pr 2 O         MoO Oi Pr 4 þ ”Zr Oi Pr 4 ” ! Zr3 Mo8 O24 Oi Pr 12 i PrOH 4

(38) (39)

It is important to underline that this approach is not applicable in the preparation of temperature-sensitive alkoxide derivatives such as those of Ni, Cu, Zn, Cd, Pb, or Bi, where the heating results in formation of metals or oxides (see section “β-Hydrogen Elimination”).

Interaction of Metal Complexes with Organic Ligands and Metal Alkoxides or Chemical Modification of Complexes in Solutions (Method 2.5) This approach provides an important alternative to the method 2.2 in the preparation of derivatives of late transition elements (the homometallic alkoxides of those being insoluble and not reactive polymeric solids). The reaction stoichiometry and conditions are dependent on the nature of reactants and on the composition of the product to be obtained. In some cases, the reaction is facile and provides the desired products as the result of mixing the reactants in proper ratio (usually in toluene) and subjecting them to short reflux, for example (Boulmaaz et al. 1997; Kessler 2003):

2

The Synthesis and Solution Stability of Alkoxide Precursors

47

    CdðOAcÞ2 þ 2Nb Oi Pr 5 ! CdNb2 ðOAcÞ2 Oi Pr 10

(40)

    CuðacacÞ2 þ 2Al Oi Pr 3 ! CuAl2 ðacacÞ2 Oi Pr 6

(41)

Preparation of the heterometallic complexes of zirconium from Zr(OiPr)4(iPrOH) requires prolonged refluxing in order to generate the reactive oxo-species (see section “Synthesis of Heterometallic Alkoxide Precursors”) (Hubert-Pfalzgraf 1994):       PbðOAcÞ2 þ Zr Oi Pr 4 i PrOH ! PbZr3 O Oi Pr 10 ðOAcÞ2

(42)

When the number of additional donor atoms of chelating ligands remains insufficient for the stabilization of the proper structure, an additional modification by chelating ligands is required (Kessler 2003; Kessler et al. 2003):     MII ðacacÞ2 þ 2Al Oi Pr 3 þ 2Hacac ! MII Al2 ðacacÞ4 Oi Pr 4 þ 2i PrOH (43) MII ¼ Co, Ni, Mg The procedure should be carried out in separate steps including mixing of the homometallic reagents, refluxing them in toluene, cooling down below the room temperature, and, only then, the addition of the necessary extra amount of acidic chelating agent (β-diketone or carboxylic acid). Introduction of an organic acid into a warm solution would result in an instant gelation (due to the condensation with released alcohols producing water in situ) instead of formation of heterometallic mixed-ligand complexes.

Heterometallic Alkoxides Formed via Formation of Heteronuclear Metal-Metal Bonds or Isomorphous Substitution No specific techniques have been elaborated for these particular cases. The majority of compounds of these two classes are obtained by complex formation between homometallic species (most often in hydrocarbon solvents) at ambient conditions or via short-term reflux (method 2.1). The formation of a metal-metal bond requires interaction of electron-rich low-valent derivatives of one metal with electron-deficient, high-valent derivatives of the other, for example (Chisholm et al. 1981; Kessler et al. 1995):       W2 Oi Pr 6 þ MoO Oi Pr 4 ! W2 MoO Oi Pr 10

(44)

ReOðOMeÞ3 þ MoOðOMeÞ4 ! ReMoO2 ðOMeÞ7

(45)

In the simplest cases, the isomorphous substitution can be achieved via mixing the isostructural but chemically different species in solution (Hubert-Pfalzgarf et al. 1978):

48

V. G. Kessler

Nb2 ðOMeÞ10 þ Ta2 ðOMeÞ10 ! 2NbTaðOMeÞ10

(46)

When the bonding parameters of two metal atoms are analogous, but the molecular aggregates observed for the homometallic species are different, the formation of heterometallic complexes, following the structure type observed for only one of two elements, can be achieved, sometimes applying the solution thermolysis (method 2.4), for example (Seisenbaeva et al. 2001b): Re2 O7 þ MoOðOMeÞ4 ! Re4x Mox O6 ðOMeÞ12

(47)

It is important to mention that the same reactants (even in the same ratio) can provide different heterometallic products if different reaction temperatures are applied. For example, the interaction between NbTa(OMe)10 and Re2O7 in toluene at room temperature provides NbTa(OMe)8(ReO4)2, while the reflux of the same reaction mixture gives Nb2Ta2O2(OMe)14(ReO4)2 (Shcheglov et al. 2002).

Solution Stability with Respect to Formation of Oxoalkoxides One of the most important requirements put on application of molecular precursors in different technological procedures, and in sol‐gel technology in particular, is the demand of stability during the application procedure itself and on storage. It is, therefore, very important to know what mechanisms can lead to the changes in the properties and molecular structures of precursors and what measures should be undertaken to support the stability of solutions and to achieve reproducibility in their application. We distinguish here between the reactions leading to transformation of alkoxide ligands resulting in formation of oxoalkoxides discussed in this part and the solvent-supported reactions of ligand redistribution (solvolysis), presented in the next one. The most important reaction pathways leading to oxoalkoxides are partial hydrolysis, oxidation by oxygen from the atmosphere and dissolved in solvents, ether and ester elimination, β-hydrogen elimination, and thermal desolvation.

Partial Hydrolysis Almost all metal alkoxides (with the exception of the kinetically rendered derivatives of precious metals and the most sterically hindered complexes) are extremely moisture sensitive. Interaction with water molecules from moist atmosphere or not properly dried solvents results in drastic changes in molecular complexity and chemical composition, for example (Ibers 1963; Bradley and Holloway 1962):


TiðOEtÞ4

 3

toluene

þ H2 O ! Ti7 O4 ðOEtÞ20

(48)

2

The Synthesis and Solution Stability of Alkoxide Precursors

Nb2 ðOEtÞ10 þ H2 O!Nb8 O10 ðOEtÞ20

49

(49)

Different hydrolysis ratios, h (number of water molecules per correspondent alkoxide formula unit, M(OR)n), provide different aggregates. For example, for the titanium ethoxide, different conditions of partial hydrolysis have also provided such aggregates as Ti8O6(OEt)20 (Day et al. 1991), Ti10O8(OEt)24 (Day et al. 1991), and Ti16O16(OEt)32 (Mosset and Galy 1988). It is not always pointed out directly, but in the complex solutions, the microhydrolysis can turn out rather selective, transforming into oxoalkoxide species only one or few of the components and changing the stoichiometry of molecular precursors. The risk of uncontrolled hydrolysis should be then eliminated as thoroughly as possible: all the operations in the preparation and weighing the samples of alkoxides are to be carried out in dry atmosphere using a Schlenk line or a dry box. The solvents dried according to reliable techniques (see Errington 1997) have to be applied. It is necessary to take into account that the water molecules can appear not only due to improper drying of the system but can even be products of different side reactions. For example, they are formed on modification of (warm) alkoxide solutions with carboxylic acids (Steunou et al. 1998): RCOOH þ R’OH ! RCOOR’ þ H2 O    Ti Oi Pr 4 þ HOAc ! Ti6 O4 Oi Pr 8 ðOAcÞ8 þ . . . 

(50) (51)

Strict control of the reaction temperature and stoichiometry (modification ratio, x, – number of modifying ligand molecules per alkoxide formula unit) is very important to insure the reproducibility of further application of such solutions.

Oxidation by Oxygen from Atmosphere or Dissolved in Solvents The alkoxide groups possessing a hydrogen atom in α-position, i.e., at the first carbon atom connected to the oxygen – all primary and secondary ones – react with oxygen in basic media, forming the products of oxidation such as carbonyl compounds and water (Turova et al. 2002): RCH2 O þ O2 ! RCHðOOHÞO ! RCHO þ OH þ . . . (52) mMðOCH2 RÞn þ OH ! Mm OðOCH2 RÞnm2 þ RCH2 O þ RCH2 OH This means that the homoleptic (alkoxide-only) derivatives of alkali, alkaline earth, and rare earth metals are very sensitive in solution to the presence of even very small traces of oxygen. The reaction is proceeding with a radical mechanism, which results in a very intensive yellowish orange (in case of high concentrations of both basic alkoxide and oxygen – even brown) coloration of solutions. The reaction speed increases with the basicity of media (alkali > alkaline earth >> rare earth elements). It is much higher in alcohols than in hydrocarbon solvents and much

50

V. G. Kessler

higher for homometallic than for the heterometallic derivatives of these elements. Really rigorous precautions can (under laboratory conditions) provide formation of the samples free from oxidation products. For the alkaline earth or rare earth elements, these are always solvates with O-donor ligands such as alcohols, THF, or dme (dimethoxyethane) (Turova et al. 2002) as desolvation itself produces the oxo-species (see below). In order to provide the samples more stable in solutions, there have been reported numerous attempts of their chemical modification using acidic ligands such as β-diketonates (Arunasalam et al. 1995) or aminoalkoxides (Poncelet et al. 1991). One of the major trends in the recently reported sol‐gel preparations of the derivatives of these elements lies in application of other organic precursors than alkoxides (β-diketonates, 2-ethylhexanoates) for the preparation of solutions or application of stable heterometallic alkoxide or heteroleptic complexes (Kessler 2003; Veith 2002; Hubert-Pfalzgraf 2003).

Ether and Ester Elimination The ether elimination reaction is a spontaneous decomposition process characteristic of, in the first hand, high-valent early transition elements, such as Mo, W, Re, Nb, and possibly Ta. The reaction mechanism involves at the first step a redistribution of electron density with a heterolytic cleavage of an O-C bond as a result. The liberated alkyl-cation is transferred to a neighboring terminal alkoxide group, forming an ether molecule: Od−_Rd+ M

O –...R+ M

OR

M=O + R2O OR

The reaction speed increases in the series of homologues Me < Et 3nm

a

b Amplification

Single crystal

Mesoscale assembly

c

Mesocrystal

Oriented Aggregation Iso-oriented crystal

Fusion

Fusion

Fig. 1 Schematic representation of classical and nonclassical crystallization. (a) Classical crystallization pathway, (b) oriented attachment of primary nanoparticles forming an iso-oriented crystal upon fusing, (c) mesocrystal formation via self-assembly of primary nanoparticles covered with organics (Reproduced with permission from reference Wohlrab et al. (2005). Copyright 2005 Wiley-VCH)

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Nonhydrolytic Chemical Routes to Metal Oxide Nanoparticles Benzyl Alcohol Route This method is straightforward and produces in good yields high-quality crystalline metal oxide nanoparticles with morphologies closely related to the starting metal reactant, without any surfactant additive. It is very versatile and may involve the reaction of benzyl alcohol with metal chlorides, alkoxides, acetates, or acetylacetonates. Depending on the reaction system, different mechanisms have been postulated (Niederberger et al. 2006a, b). Metal halide-benzyl alcohol system offers the advantage of a low operating temperature (typically between 40  C and 120  C), but residual halide impurities in the final oxide materials may be detrimental for applications such as catalysis or gas sensing. The solvothermal reaction of metal alkoxides in an autoclave at 180–250  C (the boiling point of benzyl alcohol is 205  C) offers an alternative halide-free route to nanocrystalline binary and ternary metal oxides. As a matter of fact, the reaction of tungsten chloride with benzyl alcohol yielded tungstite nanoplatelets (Polleux et al. 2005), whereas the reaction of tungsten isopropoxide yielded tungsten oxide nanowires self-assembled into bundles (Polleux et al. 2006). The uniform distance between the nanowires indicated that the attachment originated from intercalated organic molecules (most probably benzaldehyde based on FTIR spectroscopy, which arose from the oxidation of benzyl alcohol). The removal of these organic molecules by adding formamide resulted in the dispersion of individual nanowires in ethanol. Taking advantage of the in-situ formation of organic-inorganic hybrid structures, Pinna and coworkers prepared lanthanide-based lamellar hybrids with outstanding emission properties, in which oxide layers were regularly separated from each other by organic layers of intercalated benzoate molecules arising from in-situ oxidation of benzyl alcohol (Ferreira et al. 2006; Karmaoui et al. 2006). Note that the reducing ability of benzyl alcohol was applied to the synthesis of pure zero-valent metal instead of metal oxide nanostructures, by reacting nickel (Jia et al. 2008) or copper (Dar et al. 2012) acetylacetonates with benzyl alcohol under thermal or microwave activation. Actually, high boiling point and high dielectric loss factor make benzyl alcohol an appropriate medium for microwave irradiation. Microwave activation is able to reduce reaction times from days to hours, making the benzyl alcohol route energy and time efficient for continuous synthesis of large quantities of products (Bilecka et al. 2008). Control of the crystal size through the initial precursor concentration and the irradiation time was demonstrated in the synthesis of various single and mixed oxide nanoparticles, such as ZnO, Fe3O4, CoO, MnO, Mn3O4, NiFe2O4, and BaTiO3 (Bilecka et al. 2011a, b; Kubli et al. 2010). Tert-Butyl Alcohol Route The benzyl alcohol route generally results in the presence of benzyl alcoholate or benzoate residues on the particles surface, which influence the physical (typically optical properties) and chemical properties of the interface (Pucci et al. 2012). The tert-butanol route offers an alternative route to nanoparticles free of strongly

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chelating surface ligands. The reactivity of tert-butanol with metal precursors can be explained by the strong inductive stabilization of an intermediate carbocation (via SN1 mechanism in Eq. 8).

Alkoxide and Ether Routes The alkoxide route was initially used for the production of TiO2 (Koo et al. 2006; Trentler et al. 1999), ZrO2 (Joo et al. 2003), or HfxZr1-xO2 nanocrystals (Tang et al. 2004) by reaction at high temperature (300  C or more) of metal chloride and alkoxide precursors in a surfactant (trioctyl phosphine oxide, oleic acid, oleylamine). The use of carboxylic acids which bind very strongly to anatase 001 provides a powerful tool for tailoring nanocrystal shape (Fig. 2) (Jun et al. 2003). Playing on the injection rate also allowed to control the phase and shape of TiO2 nanorods (Koo et al. 2006). The ether route was shown to be suitable for the synthesis of amorphous silicabased nanoparticles (SiO2, SiO2-TiO2) as well as crystalline metal oxide nanoparticles (TiO2, SnO2) (Aboulaich et al. 2009, 2010, 2011). The syntheses were performed in dilute CH2Cl2 solutions, using chloride precursors and a stoichiometric amount of iPr2O, at mild temperatures (80–150  C), in the absence of any surfactant or coordinating solvent. The reactions conditions are thus quite different from those previously reported, which involve much higher temperatures and surfactant molecules (alkoxide route) or are performed in the alcohol acting as both an oxygen donor and a coordinating solvent (benzyl alcohol route).

Fig. 2 HRTEM analyses and simulated three-dimensional shape of TiO2 nanocrystals prepared from TiCl4 and Ti(OiPr)4 in the presence of TOPO and lauric acid: (a) bullet, (b) diamond, (c) short rod, (d) long rod, and (e) branched rod. Scale bar 3 nm (Reproduced with permission from reference Jun et al. (2003). Copyright 2003 American Chemical Society)

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Carboxylic Acid Route The reaction of titanium isoproxide in oleic acid at 270  C generated TiO2 nanorods (Joo et al. 2005). The diameter of the nanorods could be controlled by adding 1-hexadecylamine as a cosurfactant. The reaction of acetic acid with titanium n-butoxide at 100  C, without any cosolvent or additive, was successfully used in the synthesis of anatase nanoparticles (Jiang et al. 2008), while a subsequent solvothermal treatment at 200  C in the presence of in-situ produced butyl acetate, followed by calcination at 400  C, resulted in the oriented aggregation of nanocrystals (see below) into spindle-shaped mesocrystals with a single-crystal-like structure (Ye et al. 2011). The obtained nanoporous anatase mesocrystals exhibited remarkable crystalline stability and improved performances as anode materials for lithium ion batteries. Acetophenone Route Ketones and aldehydes have been shown to be able to act as alternative oxygen donors. Typically, highly crystalline anatase nanoparticles of 5–20 nm in size were obtained by reacting titanium tetraisopropoxide with common ketones and aldehydes under solvothermal conditions (Garnweitner et al. 2005). Niederberger, Kessler, and Rivas groups developed an acetophenone route from metal alkoxide or acetylacetonate precursors, which was successfully applied to the synthesis of perovskite nanophosphors and magnetic ferrite spinel nanoparticles (Abtmeyer et al. 2014; Pazik et al. 2009, 2010, 2013; Vazquez-Vazquez et al. 2008, 2011; Zhou et al. 2007). Benzylamine Route Nonhydrolytic sol-gel reactions can be carried out in an organic solvent that do not act as oxygen donors but as coordinating agents, providing control over the size, shape, and surface groups of the nanoparticles (Niederberger et al. 2006a). In those cases, the oxygen is supplied by the molecular precursors themselves. The thermal reaction of metal acetylacetonates in benzylamine belongs to this category, even though benzylamine takes part as a nucleophilic reagent. Actually, a complex mechanism was demonstrated, which involved C–C bond cleavage of the acetylacetonate ligand, followed by ketimine and aldol-like condensation reactions (Pinna et al. 2005). This route was successfully used to synthesize nanocrystalline simple oxides (e.g., aluminum, gallium, indium, iron, zinc oxides) as well as a mixed oxide (ZnGaO4) (Cao et al. 2007; Pinna et al. 2005; Zhou et al. 2007). The simple one-pot solvothermal reaction of titanium isopropoxide in benzylamine led to highly ordered hybrid structures (Garnweitner et al. 2008). These structures consisted of anatase nanoplatelets that were stacked in a lamellar fashion with a small organic layer of benzylamine molecules in between. Thus, benzylamine is involved both in the reaction mechanism leading to the transformation of the titanium isopropoxide into anatase and in the shape control of the crystals by selective capping of the (001) crystal face, which favors their growth in the [001] direction and leads to the formation of nanoplatelets. Then, the benzylamine

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molecules bound to the (001) surfaces interact with each other through pi-pi interactions, driving the stacking of the nanoplatelets into highly ordered lamellar superstructures.

Oriented Attachment and Mesocrystals Oriented attachment mechanism refers to the self-organization of nanoparticles into oriented assemblies (Niederberger and Cölfen 2006). Among these ordered nanoparticle superstructures, mesocrystals, which are assemblies of crystallographically oriented nanocrystals, differ from nanocrystal superlattices, which are periodic arrangements of nanocrystals, irrespective of their mutual orientation (Cölfen and Antonietti 2005). Different examples have been given in the previous sections yet. A further illustration of the influence of organics (solvent and additional coordinating agents) not only on the size and shape of nanoparticles but also on their self-assembling behavior is provided by the synthesis of tungstite nanostructures by benzyl alcohol route (Fig. 3). The reaction of tungsten chloride without any additive yielded tungstite nanoplatelets with a relatively broad size distribution of 30–100 nm. A comprehensive study combining ex situ and in situ techniques (Olliges-Stadler et al. 2013) evidenced a nonclassical crystallization pathway, which involved the formation of polydisperse spherical particles, their arrangement into rod-like assemblies, internal reorganization into stacked platelets, and exfoliation into shorter stacks and individual platelets. Addition of the bioligand deferoxamine mesylate, prone to providing intermolecular amide-amide interactions similar to proteins, led to nanowire bundles. These nanowires were single-crystalline and exhibited a uniform diameter of 1.3 nm. On the other hand, addition of a small amount of 4-tert-butylcatechol led to anisotropic rod-like architectures with diameters between 35 and 40 nm. The rods consisted of a highly ordered stacking of organic–inorganic hybrid nanoplatelets. However, the reaction between tungsten chloride and 4-tert-butylbenzyl alcohol (which differs from benzyl alcohol only in the presence of a bulky group in para position on the aromatic ring) resulted in the formation of ribbon-like structures consisting of laterally assembled columns of stacked nanoplatelets about 1 nm thick and 4 nm width. Another example of nonclassical crystallization mechanism is given by the crystallization pathway of indium-tin-oxide nanoparticles during solvothermal synthesis in benzyl alcohol (Ba et al. 2007). A two-step process is involved. First, a sheet-like superstructures formed, consisting of an organic matrix in which small crystallites (3–6 nm) are aligned but without any crystallographic orientation. It is assumed that the organic matrix stabilizes (both kinetically and thermodynamically) the intermediate nanoparticles. However, when the nanoparticles reach a certain size, they lost their organic protection and undergo a sudden phase transformation into the final bixbyite structure.

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Fig. 3 Time-dependent TEM and HRTEM images of the tungstite nanostructures after different reaction times of (a) and (b) 10 min, (c) and (d) 20 min, (e) 60 min and (f) 240 min (Reprinted with permission from reference Olliges-Stadler et al. (2013). Copyright 2013 The Royal Society of Chemistry)

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Processing of Metal Oxide Nanoparticles Arising from Nonhydrolytic Routes Metal oxide nanoparticles arising from nonhydrolytic routes are suitable for numerous applications, not only because of the specific organophilic properties (typically as nanofillers in polymer composites) but also owing to the new collective properties arising from their macroscopic self-assembly as films or patterned coatings.

Nanofillers Polymer nanocomposites, which intimately associate organic polymers and inorganic nanofillers, have attracted considerable attention because of their high mechanical performances and their transparency. However, oxide nanoparticles that arise from flame pyrolysis or aqueous syntheses cannot be easily dispersed in organic media, due to their hydrophilic character. These processing issues can be circumvented by using oxide nanoparticles prepared by nonhydrolytic routes. Thus, monodisperse highly crystalline ZrO2 nanoparticles were prepared by benzyl alcohol route from zirconium isopropoxide isopropanol complex as a precursor, then subjected to a simple postsynthesis treatment consisting in stirring in organic solutions of fatty-acid stabilizers at room temperature. It was shown that low amounts of these stabilizers, resulting in zirconia nanoparticles containing less than 25 wt% of organics, led to completely transparent dispersions in organic media, enabling their subsequent transfer into photo-polymerisable organic monomer phases. Transmission holographic gratings based on the prepared nanocomposites demonstrated an outstanding refractive index contrast (Garnweitner et al. 2007). Polymer nanocomposites such as epoxy and polymethacrylate (PMMA) resins were also prepared by using TiO2 nanofillers arising from benzyl alcohol and tertbutanol routes (Koziej et al. 2009; Morselli et al. 2012a, b). Recently, epoxy nanocomposites were prepared from suspensions of magnetite nanocrystals arising from benzyl alcohol route. Above blocking temperature, the magnetite nanoparticles dispersed in the epoxy resin gave rise to an interacting superparamagnetic system (Sciancalepore et al. 2015). It is worth mentioning that the ether route was shown to be suitable for the synthesis of amorphous silica-based nanoparticles (SiO2, SiO2-TiO2) (Aboulaich et al. 2009) as well as crystalline metal oxide nanoparticles (TiO2, SnO2) (Aboulaich et al. 2010, 2011) in the absence of any surfactant or coordinating solvent. Typically, dispersions of amorphous silica-based nanoparticles in CH2Cl2 were stable for months at room temperature in the absence of water and could be dispersed in an hydrophobic organic polymer without any further surface modification treatment (Aboulaich et al. 2009). Evaporation-Induced Self-Assembly Typically, evaporated-induced self-assembly (EISA) process can be implemented in the presence of surfactant-stabilized nanoparticle dispersions. As a matter of fact, ultrasmall (around 3 nm size) and highly soluble anatase nanoparticles were

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synthesized from TiCl4 using tert-butyl alcohol under microwave heating, with reaction times of less than 1 h at temperatures as low as 50  C (Szeifert et al. 2010). Mesoporous titania coatings were obtained in a one-pot procedure using sols synthesized in the presence of commercial Pluronic surfactants. These coatings could be converted into anatase upon calcination at 450  C, due to a seeding effect of the previously formed crystalline nanoparticles. The high surface-to-bulk ratio of the nanocrystals and the easily accessible mesoporous structures with extremely thin walls led to a drastic acceleration of the electrochemical Li insertion and showed high maximum capacitance in Li-ion batteries. Mesoporous materials with large mesopores of about 20 nm ordered in a cubiclike arrangement were obtained via the block-copolymer-assisted assembly of crystalline tin oxide nanoparticles arising from the reaction of SnCl4 with benzyl alcohol and redispersed in tetrahydrofuran (Ba et al. 2005). In another work, combining diblock copolymer micellar lithography with the benzyl alcohol route enabled the fabrication of patterned arrays made of quasi-hexagonally organized TiO2 nanoparticles or parallel nanowires (Polleux et al. 2011). EISA processes without any surfactant additive were successfully applied to suspensions arising from nonhydrolytic reactions. Typically, the nanoparticles prepared by the ether route, which were terminated by isopropoxide and chloride groups instead of hydroxyl groups, could be concentrated and redispersed in organic solvent; they were found to bind strongly to hydroxylated surfaces, leading to the self-limiting deposition of monolayers (in the absence of water) (Aboulaich et al. 2009). The above mentioned tungsten oxide nanobundles prepared by benzyl alcohol route from tungsten alkoxide could be dispersed in ethanol and deposited onto alumina substrates by drop-coating; a subsequent calcination step at 500  C in air removed the organics from between the nanowires without changing the macroscopic fibrous morphology (Polleux et al. 2006). The resulting highly porous coating demonstrated promising gas-sensing properties.

Microwave-Assisted Deposition One great potential of microwave heating is the possibility to implement both the synthesis of nanoparticles and their deposition on various supports (Bilecka et al. 2011a). Moreover, it is possible to selectively activate the surface of a substrate with suitable microwave-absorbing properties and the deposition of one material on top of another one (Bilecka and Niederberger 2010a). Typically, the reaction of metal acetate or acetylacetonate precursors with benzyl alcohol under microwave irradiation in the presence of immersed flat or curved glass substrates resulted in the deposition of homogeneous metal ferrite films (Bilecka et al. 2011a). The microwave-assisted benzyl alcohol route was also successfully applied to Zn (II) acetate mixed with Al(III) isopropoxide or Sn(IV) tert-butoxide, yielding Al: ZnO and Sn:ZnO nanoparticles with different doping levels (Luo et al. 2013a, b). The nanoparticle dispersions could be subsequently processed into dense transparent conducting films by repeated dip-coating and microwave-assisted densification steps before annealing under air or nitrogen atmosphere.

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Nonhydrolytic Synthesis of Mesoporous (Mixed) Oxides and Their Application as Catalytic Systems Mesoporous oxides and mixed oxides are used in a wide range of applications, such as catalysis, photocatalysis, sorption, sensing, or energy storage to cite a few. Using conventional sol-gel process, the simultaneous control of the composition, structure, and texture may be problematic, or require elaborate synthetic procedures such as prehydrolysis, chemical modification, templating, or supercritical drying. Conversely, nonhydrolytic routes (particularly the ether and alkoxide routes) can provide simple and effective methods to prepare mesoporous oxide and mixed oxide xerogels. In most cases, the use of a structure-directing agent or of supercritical drying is not necessary. As in conventional nontemplated sol-gel processes, the porosity of the xerogel (the gel dried by evaporation) results from the removal of the liquid phase (solvent plus byproducts) from the gel, as far as the gel does not completely collapse under the capillary stresses that develop during the evaporation of the liquid phase. In nonhydrolytic processes, the formation of mesoporous xerogels, with sometimes outstanding textures, has been ascribed to several factors. First, the degree of condensation of the gels can be particularly high, leading to tough solid networks, able to withstand the capillary stresses. In addition, the surface tension of the liquid phase is low compared to water, and its interaction with the nonhydroxylated pore surface is weaker. However, in the absence of a templating agent the pore structure is disordered, and the texture depends on the crystallinity of the oxide, going from interconnected pores (sponge-like texture) for amorphous materials to interparticle porosity for nanocrystalline materials (Fig. 4).

Fig. 4 (a) TEM image of an amorphous SiO2-Al2O3-MoO3 mixed oxide, after (Debecker et al. 2009); (b) SEM image of a nanocrystalline mixed TiO2-V2O5 oxide, after (Debecker et al. 2010b)

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Nonhydrolytic Routes to Mesoporous (Mixed) Oxides Alkoxide and Ether Route Mesoporous xerogels can be obtained by the strictly aprotic alkoxide and ether routes in the absence of any structure-directing agent and without using supercritical drying. After calcination, the nonordered mesoporous mixed oxides obtained by these routes maintained specific surface areas and pore volumes similar to (or higher than) those reported for solids derived from (ordered) mesoporous xerogels or aerogels prepared by conventional sol-gel (Fig. 3). A wide range of oxides and mixed oxide systems have been investigated, including SiO2, TiO2, Al2O3, WO3, SiO2-ZrO2, SiO2-TiO2, SiO2-Al2O3, TiO2-Al2O3, TiO2-V2O5, SiO2-Al2O3-MoO3, Al2O3-Ag2O-Nb2O5 (Debecker and Mutin 2012; Mutin and Vioux 2009). The textural properties of the solids depend on the crystallinity of the materials and on the reaction parameters. Thus, the porosity of amorphous SiO2-TiO2 xerogels made by these routes was found to depend on the degree of condensation of the gel and on the volume fraction of liquid phase (solvent and iPrCl by-product) in the gel; accordingly, the texture of the xerogels could be tuned simply by changing the reaction time, the reaction temperature, or the volume of solvent (Lafond et al. 2004). Optimization of these parameters led to silica–titania xerogels exhibiting outstanding textures, with specific surface areas as high as 1200 m2 g1 and pore volumes up to 2.4 cm3 g1 after calcination (Cojocariu et al. 2010). Alumina xerogels prepared by the ether or the alkoxide route led to amorphous Al2O3 with very high specific surface area (up to 400 m2 g1) after calcination at 650  C; mesoporous γ-alumina with specific surface area up to 220 m2 g1 was obtained after calcination at 850  C (Acosta et al. 1994). Mesoporous crystalline TiO2-V2O5 catalysts with specific surface areas up to 90 m2 g1 after calcination at 500  C were prepared by the ether route. The narrow mesopore distribution found in these crystalline materials results from the interspace between aggregated well-calibrated nanocrystals, as shown by scanning electron microscopy (SEM) images (Debecker et al. 2010a; Mutin et al. 2006) (Fig. 4b). Interestingly, hybrid SiO2-TiO2-MeSiO1.5 or SiO2-TiO2-Me3SiO0.5 xerogels with outstanding mesoporous textures compared to conventional sol-gel were readily obtained in one step by the ether route, from the reaction of iPr2O with SiCl4 and MeSiCl3 or Me3SiCl (Lorret et al. 2006). The incorporation of organic groups offers further control on the properties of the xerogels (in this case the hydrophobicity, which led to increased catalytic performances). Acetamide Elimination Route Quite recently, Pinkas and coworkers in Brno investigated the synthesis of SiO2TiO2 and SiO2-ZrO2 mixed oxides by a nonhydrolytic route involving acetamide elimination from silicon acetate, Si(OAc)4, and titanium or zirconium diethylamide, M(NEt2)4, at 80  C (Skoda et al. 2015a, b). In these cases, the addition of Pluronic P123 as a structure-directing agent was needed to obtain after calcination homogeneous mixed oxides with a “wormhole” mesopore structure and high specific surface area. Interestingly, when an excess of silicon acetate was used to reach Si:M ratios

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higher than 1, the formation of acetic anydride was detected, suggesting that homocondensation of excess Si–OAc groups took place, probably catalyzed by metal Lewis acidic species.

Ester Elimination Route Pinkas and coworkers recently described the nonhydrolytic synthesis of silicon orthophosphate xerogels by elimination of trimethylsilylacetate, AcOSiMe3, from silicon acetate, Si(OAc)4, and the tris(trimethylsilyl)phosphate, PO(OSiMe3)3 (Styskalik et al. 2014, 2015b). The nonhydrolytic polycondensation led to the formation of an inorganic network with a high content of Si–O–P bonds and hexacoordinated SiO6 moieties. In the absence of structure-directing agent, the silicophosphate xerogels exhibited significant specific surface area but were largely microporous (Styskalik et al. 2014). When a Pluronic P123 template was added (Styskalik et al. 2015b), mesoporous silicophosphate materials with specific surface area up to 128 m2 g1 and pore diameters around 20 nm were obtained after calcination. This method has been extended to the synthesis of hybrid “silicophosphonate,” starting from organotriacetoxysilanes, RSi(OAc)3, with trimethylsilyl esters of phosphonic acid, RPO(OSiMe3)2 (Styskalik et al. 2015a). R2Si(OAc)2 precursors were not incorporated in the network but acted as templates. Bridged acetoxysilanes, (AcO)3Si-X-Si(OAc)3, and bridged trimethylsilylphosphonates, (SiMe3O)2PO-XPO(OSiMe3)2, have also been used, leading to hybrids with large specific surface area (up to 700 m2 g1) and pore volumes (up to 1.6 cm3 g1). Interestingly, the microporosity of these bridged hybrid materials is correlated to the presence of hexacoordinated SiO6 structural units. Carboxylic Acid Route The reaction of silicon alkoxides with formic acid, originally reported by Sharp (1994), leads to low porosity silicas. However, mesoporous silicas with a surface area of 720 m2 g1 and pore volume of 1.4 cm3 g1 could be obtained via the formic acid route by using an ionic liquid as solvent and long aging times. In this case, the ionic liquid was then washed off with a polar solvent (Dai et al. 2000). This route was also applied to the synthesis of crack-free ionogel monoliths, consisting of an ionic liquid phase confined in a mesoporous silica (Neouze et al. 2006; Viau et al. 2012). Ionogels are well suited for the in situ fabrication of temperature-resistant electrolytes for energy storage devices, including thin film electrochemical double layer capacitors (Horowitz and Panzer 2012).

Application to Heterogeneous Catalysis The mesoporous oxides and mixed oxides prepared by nonhydrolytic routes have been successfully applied to the design of heterogeneous catalysts and photocatalysts, as recently reviewed (Debecker et al. 2013; Debecker and Mutin 2012).

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In particular, the ether route provides a simple, one-step method to prepare mesoporous mixed oxide catalysts with well-controlled compositions and textures, starting from cheap chloride precursors and avoiding the use of reactivity modifier, templating agent, or supercritical drying step (Debecker et al. 2013). When solubilization of a chloride precursor is problematic, the alkoxide route may be more indicated (Helena Kaper et al. 2012). The nonordered texture of the catalysts prepared by these routes is well-suited to catalytic applications. The large, interconnected mesopores featured by these materials are more favorable than ordered mesopores. These routes have been used for the synthesis of a wide variety of nonordered mesoporous catalysts (Debecker et al. 2013), such as SiO2-TiO2 (Cojocariu et al. 2008), SiO2-ZrO2 (Helena Kaper et al. 2012), SiO2-WO3 (Maksasithorn et al. 2014), TiO2-V2O5 (Mutin et al. 2006), Ag-Nb2O5-Al2O3 (Petitto et al. 2013), SiO2-Al2O3-MoO3 (Debecker et al. 2009), and SiO2-Al2O3Re2O7 (Bouchmella et al. 2013), which were tested in various reactions, including mild and total oxidation, alkylation, selective catalytic reduction of NOx by NH3 or decane, and olefin metathesis. In the case of mixed oxides, the catalytic properties strongly depend on the degree of homogeneity. Nonhydrolytic methods often lead to highly homogeneous mixed oxide xerogels, owing to their easily controllable kinetics and, in the case of metal silicates, to the levelling of reactivities at silicon and metal centers. When the oxide components feature high Tammann temperatures, as for instance in the case of SiO2TiO2, SiO2-ZrO2, or SiO2-Al2O3, homogeneity is maintained after calcination (typically at 500  C). For instance, SiO2-TiO2 mixed oxides obtained by the ether route show excellent performances in the mild oxidation of organic compounds, ascribed to the formation of well-dispersed Ti species linked to the SiO2 matrix by Si–O–Ti bonds and to their very high specific surface area (Cojocariu et al. 2010; Lafond et al. 2002). On the other hand, when the active oxide phase has a low Tammann temperature (cases of VOx, MoOx, ReOx, WOx for instance), appropriate thermal treatments can provoke the migration of active oxide species toward the surface, leading to an increase of their surface concentration (Debecker et al. 2013). For instance, highly active MoO3-SiO2-Al2O3 and Re2O7-SiO2-Al2O3 olefin metathesis catalysts have been obtained by the ether route (Bouchmella et al. 2013; Debecker et al. 2009). Migration of Mo or Re oxide species toward the surface occurred during the calcination at 500 C, as evidenced by XPS and ToF-SIMS, leading to high concentration of well-dispersed Mo or Re surface species. This, together with the mesoporous texture and acidic character of the catalysts, accounted for their excellent catalytic performances in the metathesis of ethene and butene to propene (Bouchmella et al. 2013; Debecker et al. 2009). Recently, mesoporous SiO2-ZrO2 obtained by the templated acetamide elimination route were found to display good activity and selectivity in the Meerwein–Ponndorf–Verley reduction of 4-tert-butylcyclohexanone and for aminolysis of styrene oxide with aniline (Skoda et al. 2015b). Mesoporous SiO2TiO2 xerogels prepared by the same route were tested in cyclohexene epoxidation with cumyl hydroperoxide in toluene. They displayed a good catalytic activity in cyclohexene epoxidation with cumyl hydroperoxide (Skoda et al. 2015a).

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Conclusion Nonhydrolytic sol-gel is now well established as a powerful methodology for the synthesis of oxide-based materials, notably mixed-oxide xerogels with wellcontrolled composition and mesoporosity, highly crystalline oxide nanoparticles with control over size and shape, as well as superstructures and films resulting from their assembly. One major asset of this low-temperature process is to offer simple synthesis protocols avoiding the use of chemical additives or templates. Actually, the control over composition, texture, structure, and morphology of the final materials arises from the intrinsic nature of the organic reactions involved. Typically, in the synthesis of mixed oxides, nonhydrolytic condensation reactions are slower than hydrolytic ones, while leading to higher condensation degrees and hydroxyl-free surfaces. As the reaction rates depend more on the nature of the carbon center than on the nature of the metal center, the reactivity of the different metal precursors is leveled. As for the reactivity of silicon precursors, which is much lower than that of metal precursors, the reactions around the silicon atom are catalyzed by Lewis acidic metal species, which again levels the kinetics when metal-silicates are prepared. In the synthesis of oxide nanoparticles, the presence in the nonaqueous reaction mixture of organic derivatives endowed with complexing ability strongly influences the size, shape, and structure of crystallites and also their assembly into superstructures. Even though the various organic reactions involved in nonhydrolytic processes may make it difficult to predict the characteristics of the final material, in practice, this complexity offers wide possibilities to tailor the reaction system, by varying the precursor-oxygen donor combination. Recently, the use of microwave chemistry has even considerably enlarged this potential. As a matter of fact, over the last decade, nonhydrolytic reactions have been extended to other processes (solvothermal synthesis, chemical solution deposition, atomic layer deposition), solvents (ionic liquids), O-donors (including cellulose), activation modes (microwaves), and materials (metal chalcogenides or phosphates, carbon nanocomposites) (Bilecka and Niederberger 2010b; Mutin and Vioux 2013). Such a development in innovative chemical methods should be continued, owing to the increasing demand for advanced nanostructured materials, for such applications as energy storage, catalysis, sensing, and optics.

References Aboulaich A, Lorret O, Boury B, Mutin PH. Surfactant-free organo-soluble silica-titania and silica nanoparticles. Chem Mater. 2009;21:2577–9. Aboulaich A, Boury B, Mutin PH. Reactive and organosoluble anatase nanoparticles by a surfactant-free nonhydrolytic synthesis. Chem Mater. 2010;22:4519–21. Aboulaich A, Boury B, Mutin PH. Reactive and organosoluble SnO2 nanoparticles by a surfactantfree non-hydrolytic sol-gel route. Eur J Inorg Chem. 2011;2011:3644–9. Abtmeyer S, Pazik R, Wiglusz RJ, Malecka M, Seisenbaeva GA, Kessler VG. Lanthanum molybdate nanoparticles from the Bradley reaction: factors influencing their composition, structure,

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Pechini Processes: An Alternate Approach of the Sol-Gel Method, Preparation, Properties, and Applications

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Lucangelo Dimesso

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Pechini’s Process and Method: Preparation and Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Pechini’s Process and Method: Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . High-Tc Superconductors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lithium Transition Metal Phosphates as Cathode Materials for High-Voltage Li-Ion Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Concluding Remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

This chapter highlights a process related to the sol-gel route called “Pechini, or liquid-mix, process.” This process overcomes most of the difficulties and disadvantages that frequently occur in the alkoxide based sol-gel process. After a short historical introduction on the inventor, chemical mechanisms of the process, the advantages on the preparation of ceramic materials, and the drawbacks of the process will be shown. Finally, as examples of applications of the Pechini method, the results of investigation on ceramic superconductors and on Li-ion containing metal phosphates as cathode materials for Li-ion batteries are presented. The versatility, extensive applicability of the Pechini process with low production costs and potential material applications are demonstrated.

L. Dimesso (*) Earth and Material Sciences Department, Technische Universitaet Darmstadt, Darmstadt, Germany e-mail: [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_123

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Introduction In our industrialized world, there is a growing need for mixed-cation oxide ceramics. Often these advanced ceramic materials have composition selected for their unique electrical or dielectric properties. Typical examples include lead zirconate-lead titanate (PZT) materials doped with a small amount of niobium for piezoelectrics, cation-doped barium titanate or strontium titanate for capacitors, lead magnesium niobate for dielectrics, yttrium barium copper oxides (YBCO) for high critical temperature (T > 77 K) superconductors, and strontium-doped lanthanum manganite for semiconducting electrodes (Lessing 1989). These compositions often contain four or more cations. After the Second World War (WWII) in the 1950s and1960s, especially for applications in an industrial scale, the conventional method of preparing powder formulations used in the manufacture of ceramic materials required the calcinations of a mechanically ground mixture of metal oxides and carbonates in definite proportions. However, the milling and grinding, normally employed to obtain a mixture in a fine state of subdivision, introduces contaminates from abrasive materials. These contaminants have a detrimental effect on the electrical properties and introduce a variance into each batch of powder prepared. The solid-state reaction, a diffusioncontrolled process, requires intimacy of reacting species and a uniform distribution of each species to obtain a completely reacted and uniform product. The mechanically ground mixture requires prolonged calcinations at high temperatures under accurate control of the atmosphere. Such prolonged calcination promotes crystallite growth which is undesirable in the fabrication of dense fine-grained ceramics. Moreover, it is extremely difficult to prepare dielectric films composed of two or more chemically combined oxides by conventional evaporation techniques. This requires the difficult art of controlling the rate of deposition of metal oxides from the vapor phase onto a substrate enclosed in an evacuated chamber. The high temperature and high vacuum required to vaporize the oxides cause variations in oxidation states and contamination from metal vapors. For these reasons a process eliminating the cumbersome apparatus and tedious techniques was needed. In order to produce mixed oxide dielectric films economically and to obtain optimum and reproducible electrical properties, it was necessary to eliminate mechanical mixing and lower calcination temperatures were required. In the same years (1950s–1960s), sol-gel process was moving the first steps, but it was not established yet as one of the most important chemical methods for the preparation of materials. Indeed Roy and Osborn (1954) and Roy and Roy (1955) recognized the potential for achieving very high level of chemical homogeneity in colloidal gels and synthesized a large number of novel ceramic oxide compositions, involving Al, Ti, Si, and Zr (transparent oxides). Under these circumstances, in order: – To provide a process for forming high purity dielectrics – To provide a method for precisely controlling the constitution of a high purity dielectric

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– To provide a method of forming a thin layer of high purity dielectric on metallic substrate and to form high purity piezoelectric compositions M. P. Pechini developed a sol-gel method for lead and alkaline-earth titanates and niobates, materials which do not have favorable hydrolysis equilibria (Pechini 1967a, b). This method is also known by the names Pechini and liquid-mix process. Maggio Paul Pechini (born on May 1, 1922, Dauphin County (Pennsylvania), dead on June 1, 2007 (85 years old), in Brandon, Hillsborough County (Florida)) worked, at that time, for the Sprague Electric Company in Massachusetts, which was interested in manufacturing ceramic capacitors. Sprague Electric was started by Robert C. Sprague in 1926 in Quincy, Massachusetts, as Sprague Specialties Company. R. C. Sprague is credited with inventing a tone control device that greatly improved the sound of radios. In 1929, the company decided to move to North Adams, Massachusetts. The move was completed in 1930. At its peak, Sprague Electric employed 12,000 people worldwide, including over 4000 people in North Adams alone in five separate sites. The largest site (a former textile mill), called “Marshall Street,” was composed of 23 different buildings, all linked by covered overpasses and tunnels. The site is now the home of Mass MoCA (Massachusetts Museum of Contemporary Art). In the mid-1960s, Sprague had plants in Scotland, France, Italy, and Japan, in addition to multiple locations in the USA (Fig. 1, adapted from the Massachussets College of Fine Arts (MCLA) library website). In 1976, Sprague Electric was acquired by General Cable (renamed GK Technologies in 1979). In 1981, Penn Central acquired all the stock of GK Technologies. In 1985, the Sprague North Adams operations were closed. At this time, Sprague was manufacturing a wide range of products including tantalum, aluminum, film, paper, and ceramic capacitors, resistor networks, pulse transformers, and filters. The major manufacturing sites for tantalum capacitors were located in Sanford (Maine), Concord (New Hampshire), and Tours, France. In 1992, Vishay acquired the tantalum capacitor operations of Sprague Electric Company, while other product lines were spun off. In the last 30 years, the original “Pechini process” has been further developed and extended to the synthesis of electric and magnetic materials, including ferroelectric and capacitor materials; superconducting, photocatalytic, magneto-optical, electrolytic materials for solid oxide fuel cells; and so on. Several other solvents and starting materials have been also used which contributed to the diffusion and success of the Pechini process. This process has been developed to a Pechini method and is applied worldwide from labor to industrial level. Even the number of the publications on this subject has dramatically increased that a summary in this chapter would be a huge challenge. Then, a short description on the improved material properties prepared by Pechini methods compared to other methods (such as solid-state reaction method and amorphous citrate method) for a very few applications, quoting the corresponding literature (particularly reviews and extended articles), can be given.

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Fig. 1 (a) Sprague Electric logos from 1926 to 1985; (b,c) Sprague purchased the Arnold Print Works plant in North Adams, MA. This became the permanent headquarters and principal manufacturing site for Sprague Specialties Company (Adapted from the Massachusetts College of Fine Arts (MCLA) library website)

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Pechini Processes: An Alternate Approach of the Sol-Gel Method. . .

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The purpose of this work is to reveal the feasibility, versatility, advantages, and disadvantages of the Pechini method for the synthesis of some materials, in an effort to gain fine control of the material morphology and find novel materials. After showing the chemistry and mechanisms of the Pechini process (and method in some cases) and the influence on the structural, microstructural, and corresponding physical properties, although many kinds of complex oxide are attempted to be prepared via Pechini methods, superconducting materials composed of various kinds of elements and lithium transition metal phosphates, powders and composites, are chosen for examples of applications.

Pechini’s Process and Method: Preparation and Properties The Pechini approach has been popularized by H. Anderson, who originally applied the method to fabricate perovskite powders for high-temperature magnetohydrodynamic electrodes in the 1970s. Eror and Anderson (1986) reported that the method has been used successfully on over 100 different mixed oxide compounds including lanthanum manganite and barium titanate. The process calls for forming a chelate between mixed cations (dissolved as salts in a water solution) with a hydroxycarboxylic acid (citric acid is preferred) (Lee et al. 2003). The cations are chelated and then, with the aid of polyalcohols, the chelates are cross-linked to create a gel through esterification. This has the distinct advantage of allowing the use of metals that do not have stable hydroxo species. Initially, Pechini used citric acid (CA). This has often been replaced with EDTA (ethylenediaminetetraacetate), which has the advantage of chelating most metals and, with four carboxylate groups, is easily cross-linked to form gel. It is also possible to use polyvinyl alcohols that provide for a three-dimensional network during gel formation as shown in Fig. 2. The gelled composites are sintered, pyrolyzing the organic and leaving

Fig. 2 Scheme of the Pechini method reactions (Reproduced from Lee et al. (2003) with permission of The Royal Society of Chemistry)

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nanoparticles, which are reduced by the pyrolyzed gel. Various cation salts can be used, such as chlorides, carbonates, hydroxides, isopropoxides, and nitrates. The general idea is to distribute the cations atomistically throughout the polymer structure. Heating (calcinations) of the resins in air or other gases causes a breakdown of the polymer and “charring” at about 400
C. It is assumed that there is little segregation of the various cations that remain trapped in the char. Subsequently, the cations are oxidized to crystallites of mixed cation oxides at 500–900
C. The process is quite complicated, and it has many changeable experimental variables that affect the final product. There are two basic chemical reactions involved in the Pechini process to make ceramic “precursors”: (1) chelation between complex cations and citric acid and (2) polyesterification of excess hydroxycarboxylic acid with glycol in a slightly acidified solution. The possible polyesterifications occurring in the liquid-mix process have been reported by Anderson’s group (1987), Weber et al. (2005), and Zhang et al. (1990). This viscous liquid is then dried, by applying heat or vacuum, to form a gelatinous precursor for ceramic powders (or films). A final calcination removes all organic substances to yield oxide powders. Additional pulverization is usually required to break down any hard agglomerates in final powders formed during charring and calcination. The limitation of the Pechini method, like many techniques, lies in the lack of size and morphological control. With traditional sol-gel methods, the particles are part of a gel structure, while in the Pechini method, the metal cations are trapped in the polymer gel. This reduces the ability to grow controlled shapes and involves the formation of hard crystallite agglomerates. The size of the final product is controlled, to an extent, by the sintering process and the initial concentration of metals in the gel. Evidence (Lessing 1989) indicated that the physical morphology of final oxide powder made by a Pechini-type method was mainly influenced by the morphology of its resin intermediate. Fine, non-agglomerated powders cannot be obtained by calcining a dense, rigid resin intermediate without extensive post calcination grinding. Till 1992, the published literature gave no answer to the question of how the chemistry of CA–ethylene glycol (EG) polymeric precursors influences the morphology of a resin intermediate (or ceramic powder) in this type of process. Tai and Lessing (1992) investigated the optimal composition in citric acid–ethylene glycol (CA–EG) polymeric mixtures. The authors claimed that with a proper composition, the subsequent polymerization, decomposition, and charring of the polymeric precursor can be simplified. Rheological properties of various CA–EG mixtures were first characterized to determine the optimal composition for resin formation. Thermal analyses on CA–EG mixtures were conducted to further verify this composition. Polymeric mixtures yielding the most foamed resin intermediates, with or without ceramic constituents, were found experimentally and compared to a result of gel-point calculation. Multicomponent ceramic powders of lanthanum chromite doped with 15 mol% SrO (LSC) were prepared during the investigation. To investigate this mechanism, Tai and Lessing (1992) provided direct comparisons on the foaming behavior of polymeric bare gels and LSC ceramic precursors made of CA and EG.

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Pechini Processes: An Alternate Approach of the Sol-Gel Method. . .

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1.0 × 107

A series of beakers containing homogeneous mixtures of CA and EG in varied ratios (with equal amounts of hot water and nitric acid) were kept warm on hot plates while stirring. Before the charring test started, all beakers were sealed by plastic film to maintain the water content. Immediately after removing the covers and stirring rods, sample beakers were placed into an oven. These highly viscous samples started softening, steaming, boiling, frothing, fuming, and finally forming dry sponges within a few minutes. A direct comparison on the volumes and macromorphology of resultant resin was then made. The foaming behavior – of LSC precursors – was then characterized following the same procedures for bare-gel samples. Viscosities of bare polymeric gels as a function of temperature are plotted in Fig. 3. At any temperature, the viscosities of CA–EG mixtures increased with increasing CA content and reached a maximum at 50 mol% CA. Note that when the CA content changed from 30 to 40 mol%, the resulting viscosity suddenly increased about two orders of magnitude. This indicates that a gelling process starts rapidly at this composition. There was no such change in viscosities of samples containing CA from 40 to 60 mol%. The high viscosity values dropped a little at CA = 65 mol% followed by a significant decline as the CA content reached 70 mol%. Citric acid did not dissolve completely in the sample containing 80 mol% CA. Figure 3 shows that the gelling reaction takes place when citric acid content falls between 40 and 60 mol%. Step-reaction polymerizations are

1.0 × 105 1.0 × 104 1.0 × 103

80/20 70/30 65/35 60/40 55/45 50/50 40/60 30/70 20/80

1.0 × 10

1

1.0 × 10

2

log Viscosity (cps)

1.0 × 106

CA / EG in mole

0.

25

50

75

100

125

Temperature (°C) Fig. 3 Viscosity at various CA–EG bare gels as a function of the temperature (Reproduced with permission of Materials Research Society)

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usually enhanced by making one of the reactants excess in amount. For a system containing CA and EG, excessive EG has been historically preferred because of its lower cost and superior solubility. These might explain why in the original process Pechini used a CA–EG mole ratio of 20/80. According to the theoretical calculation, gelation should start when more than 40 mol% CA is added into EG or when more than 43 mol% EG is mixed into a bath of CA. Actually, the esterification reaction took place in all compositions of CA–EG mixtures, and it ceased when one of the reactants was consumed in that mixture. For those samples containing excessive CA or EG, the polymeric gel was quickly diluted by the “parent solution” and caused low viscosities. Nevertheless, a less-likely possibility of forming other types of ester in the studied system cannot be excluded. Foaming characteristics of various bare gels were directly compared by the apparent volume of each charred resin. Upon heating, both gel samples of CA/EG = 60/40 and 65/35 were frothing continuously and smoothly. The resultant resins had the largest volumes among all samples. Fierce bubbling and fuming were seen in those beakers containing either superfluous CA or EG. Furthermore, the foaming process of those samples with excessive CA or EG was interrupted by a tendency to shrink before they solidified. The foaming behavior of CA–EG bare mixtures was somewhat different from the results of viscosity measurements (which suggested a well-gelled resin to form in the sample with CA mole% = 40–60). This phenomenon can be explained by the fact that all sample beakers were quickly heated up in a closed oven. Therefore, excessive glycol within those samples rich in EG did not evaporate out before a resin formed. Such less volatile EG either filled in the open structure of a semi-dried resin or glued to the surface of the polymeric skeleton during the foaming process. Consequently, the expansion of a polymeric gel upon heating was retarded by the sticky EG residues. The remainder of EG in the beaker could not be burned off at charring temperatures (200–300
C), and it solidified by contraction when the resin temperature decreased. The authors observed significant shrinkage of the polymeric sponge during cooling in the sample with high glycol content. It can be concluded that excess EG is the major barrier to the foaming of CA–EG resins in Pechini-type processes. On the other hand, any free CA in a polymeric gel might result in complex reactions due to possible stepwise decompositions of citric acid, such as Cadogan, MacWilliam et al. (Dictionary of organic compounds 1965) citric acid ! aconitic acid ðC6 H6 O6 Þ þ H2 OðgÞ; aconitic acid ðheatÞ ! itaconic acid ðC5 H6 O4 Þ þ CO2 ðgÞ; itaconic acid ðheatÞ ! itaconic anhydride ðC5 H4 O3 Þ
þ H2 OðgÞ, or acetonedicarboxylic acid COðCH2 • COOHÞ2 : The above derivatives from CA are also carboxylic acids which can react with metal nitrates or glycol in the mixture. It thus suggests that an excess amount of CA probably will complicate the esterification reactions occurring in the CA–EG

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Pechini Processes: An Alternate Approach of the Sol-Gel Method. . .

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Fig. 4 (a) Charred bare resins with various ratios of CA–EG; (b) charred LSC resins with various ratios of CA–EG (Reproduced with permission of Materials Research Society)

polymeric system by forming other polymeric species. As a result, the foaming behavior of a polymeric gel rich in CA will be difficult to predict and control. When an aqueous LSC precursor was prepared by adding a premixed polymeric (CA–EG) gel into the nitrate solution on a hot plate, there was no precipitation observed throughout the whole process. Upon heating, brown fumes were first released from the mixture of LSC nitrate and polymeric gel when it reached about 75
C. The viscosity of this solution did not change much after this strenuous reaction; i.e., no apparent sign of polymerization occurred at this stage. Frothing of each mixture started right after a strenuous steaming occurred above 110–120
C, and finally dry resin formed at about 250
C. The above observations indicated that nitrates decomposed at lower temperatures than citrates, as previously suggested by Tuttle and Voigt (1988) for a similar powder process. When all ingredients were directly mixed in a beaker, the above reactions took place in an unpredictable manner. Figure 4b shows the foaming properties of various LSC charred resins containing CA–EG polymeric gel. The most expanded resins were made from LSC gels with 40–55 mol% CA. This composition range is identical to the calculated gel composition (40–57 mol%). By comparing Fig. 4a, b, the foaming nature of CA–EG bare gel evidently contributed to the formation of porous ceramic precursors in the Pechini-type process, although the morphological dependence on gel composition was slightly different between bare polymeric gels and LSC gels. The appearance of the LSC resin is similar to that of a charred sponge made of the same polymeric gel. The sticky mass left on the inside walls of the oven used for resin setting indicated that significant amounts (but not all) of the glycol had evaporated during heating. The removal of excess reactant EG by steam during the thermal process might be responsible for the changes in foaming behavior of the CA–EG polymeric less-acid regions. Therefore, the introduction of a premixed solution of CA and EG at proper ratios, instead of pure CA powder, into the nitrate solution should be beneficial. The removal of excess reactant EG by steam during the thermal process might be responsible for the changes in foaming behavior of the CA–EG polymeric gels previously mentioned. A glycol-rich polymeric gel (80 mol% EG) was used in the original Pechini approach, which usually caused dense agglomerates of resin intermediate. It can be understood by the fact that large quantities of excess glycol were

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not removed from the polymeric precursor at the charring step. The LSC charred resins were either dense flakes or large lumps when a CA-rich or EG-rich gel was used, respectively. The empirical gel composition for making a heavily expanded LSC resin was found at CA–EG = 50/50. These ceramic semi-products were amorphous, as analyzed by x-ray diffraction. DTA curves were measured to reveal the thermal decomposition behavior of each mixture. Examination of the portions of the DTA curves between 190
C and 250
C revealed no detectable endothermic evaporation of excessive EG from the CA–EG samples at the boiling temperature of EG. It indicates that there was a large quantity of EG (or its derivatives) confined in the polymeric precursor even though the resin was charred. As discussed earlier, this residual glycol is detrimental for the resin to expand its volume. These DTA results support the previous explanation made on the foaming–shrinking behavior of an EG-excess sample. It was also found that residual glycol made the thermal decomposition of charred resins intricate. Since all bare resins were already formed between 250 and 300
C, further thermal analyses of these samples up to 800
C gave information of the carbonization process in charred polymers. In the DTA profiles (Tai and Lessing 1992), two main exothermic peaks are seen at about 390
C and about 560
C, respectively. When the sample composition falls within the theoretical gelling region (40–57 mol% CA), only one exothermic peak is observed. A new peak appears between 480 and 570
C in those samples having EG content less than 35 mol%. This new exothermic reaction could be correlated to ignition of pure CA. These DTA results demonstrate that a simple, continuous polymerization and decomposition occurred in the CA–EG polymeric gel containing about 55–65 mol% citric acid. LSC “liquid precursors” with various gel compositions, were heated from room temperature to the calcination temperature directly in situ. The appearance of split or single peaks of varied precursors is similar to the phenomena occurring in samples of bare gel. A distinct, strong peak was observed in the LSC sample having CA content between 40 and 60 mol%. The ending temperature of this exothermal reaction in most samples was around 410
C, which was about 200
C lower than the ending temperatures of corresponding CA–EG polymeric gels (610
C). Because no significant exothermic peak was detected in each run beyond 500
C, it can be concluded that most of the organic substance in LSC precursors has been burned out in flowing air by this temperature. An interesting phenomenon was observed when a homogeneous mixture of polymeric gel and LSC nitrate solution was set in a “big evaporating dish” and stored in a drying oven at about 35
C. After several days (as water evaporated), a huge, porous, dry sponge of LSC resin formed whose appearance and microstructure were identical to a resin with the same composition but charred at 250
C. The removal of water and other gaseous species from wet gels also can be achieved by applying vacuum. Since this low-temperature foaming action took place significantly only in those LSC precursors containing equimolar CA and EG, it provided another evidence to support the suggested gelling composition in the studied system. Without more detailed information on the fundamental chemical bonding configurations, it is impossible to unequivocally identify the exact chelating reactions between the polymeric gel and metal ions in the studied system. However, there is

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sufficient evidence to show that a highly porous polymeric resin (most likely polyester) is formed by dehydrating an acidified mixture of citric acid and ethylene glycol in proper ratio. This organic foam is an ideal skeleton where chelating reactions can take place homogeneously in the Pechini-type powder processes.

Pechini’s Process and Method: Applications This part of the chapter is dedicated to some important applications of the Pechini process. Advanced ceramic materials have compositions selected for their unique electrical or dielectrical properties. The method was first developed by Pechini to prepare capacitor materials focusing only on niobates, titanates, and zirconates. After Pechini this method has been extensively applied toward the synthesis of a variety of multicomponent oxides (see Table 1, Kakihana and Yoshimura 1999). All the studies reported have clearly indicated that the Pechini method is quite suitable for producing highly pure and homogeneous oxides at reduced temperatures (400–900
C). Moreover, the Pechini process has been used for the synthesis of electric and magnetic materials rather extensively, including ferroelectric and capacitor materials, superconducting materials, photocatalytic materials, magneto-optical materials, and electrolytic materials for solid oxide fuel cells (Kakihana 1996; Reichenbach et al. 2003; Yoshino et al. 2002). The improved material properties for the PSG process with respect to the other methods (such as solid-state reaction method and amorphous citrate method) have been demonstrated by Kakihana Table 1 List of multicomponent oxides prepared by Pechini method. The corresponding references are labeled as in the source. (Source: Kakihana and Yoshimura 1999)

Materials Lead magnesium niobates LaMnO3 LaAlO3 Sr-doped lanthanum chromite Y3NbO7 LiTaO3 K2La2Ti3O10 ZrO2/CeO2 ZrO2/Y2O3 ZrO2/Y6WO12 BaSnO3 BaTi4O9 KTiNbO5 SrTiO3 BaTiO3 PbTiO3 La2Ti2O7 Y2Ti2O7 Cuprates

Reference numbers 15 14, 16, 50 51 52 53 54 55, 56 57 58 59, 60 61 63 64 17–22 23–27 28 30 29 31–49

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(Kakihana and Yoshimura 1999) in a review article for the synthesis of superconductors and photocatalysts. On the other hand, Lin et al. (2007) have applied the Pechini process to the systematic synthesis of various kinds of oxide optical materials, mainly luminescence and pigment materials with different forms (powder, core–shell structures, thin film, and patterning). The purpose of the research was to reveal the feasibility, versatility, advantages, and disadvantages of this method for the synthesis of such optical materials, in an effort to gain fine control of the material morphology and find novel optical materials. In the feature article, the authors demonstrated the multiform of the optical materials derived from the Pechini precursor solutions, including powder luminescent materials (combined with the spray drying process), monodisperse and spherical core–shell structured phosphor and pigment particles (via the surface modification process), thin-film phosphors (via dip coating), and their patterning (combined with the soft-lithography process). Due to the very wide range of applications of the Pechini process, to deal with all of them in this chapter would result in a huge effort. For this reason and on the basis of my research activity, I will focus my attention on examples of application on superconducting ceramics (called also “high-Tc superconductors”) and on Li-ions containing transition metal phosphates used as cathode materials for high-voltage Li-ion batteries. As far as other applications, the recommendation to our readers is to use our references as sources for their research.

High-Tc Superconductors High-temperature superconductors (high Tc or HTS) are materials that behave as superconductors at unusually high temperatures. The first high-Tc superconductor was discovered in 1986 by IBM researchers G. Bednorz and K. A. M€uller (Bednorz and M€ uller 1986) who were awarded the 1987 Nobel Prize in Physics “for their important breakthrough in the discovery of superconductivity in ceramic materials.” Whereas “ordinary” or metallic superconductors usually have transition temperatures (temperatures below which they superconduct) below 30 K (243.2
C) and must be cooled using liquid helium in order to achieve superconductivity, HTS have been observed with transition temperatures as high as 138 K (135
C) and can be cooled to superconductivity using liquid nitrogen. Until 2008, only certain compounds of copper and oxygen (so-called cuprates) were believed to have HTS properties, and the term high-temperature superconductor was used interchangeably with cuprate superconductor for compounds such as bismuth strontium calcium copper oxide (BSCCO) and yttrium barium copper oxide (YBCO). The simplest method for preparing high-Tc superconductors is a solid-state thermochemical reaction involving mixing, calcination, and sintering. The appropriate amounts of precursor powders, usually oxides and carbonates, are thoroughly mixed using a ball mill. Solution chemistry processes such as coprecipitation, freezedrying, and sol-gel methods are alternative ways for preparing a homogeneous mixture. These powders are calcined in the temperature range from 800
C to 950
C for several hours. The powders are cooled, reground, and calcined again.

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This process is repeated several times to get homogeneous material. The powders are subsequently compacted to pellets and sintered. Although coprecipitation and sol-gel techniques were developed (Zhang et al. 1990) based on sophisticated organometallics for decreasing the diffusion path, syntheses became even more difficult with the multiphase systems of bismuth and thallium. Multiphase families of bismuth and thallium demanded more accurate control for correct stoichiometry because of the copper complexation and the remarkable solubility product distinctions. In order to improve the chemical quality of the superconducting powders, and consequently the physical properties, Chiang et al. (1993) modified the Pechini method as reported in the flowchart in Fig. 5. Metal cations were provided by corresponding salts of nitrate, except for the thallium. They were complexed by a bidentate ligand supplier, namely, oxalic acid. Oxalic acid was favored because of the faster reaction rate than citric acid, a multidentate ligand. By appropriate selection of the chelating agent, the synthetic approach was demonstrated to overcome most of the problems encountered in other synthetic routes. Multiphase compound was a major beneficiary because of the straightforward phase refinement. The authors concluded that the Pechini method is considered valuable for mass production of high-quality superconducting powders. In another example, Peng et al. (1998) adapted the Pechini process to the synthesis of Bi-2223 ((Bi,Pb)2Sr2Ca2Cu3O6) system and its various modified forms. The Pechini process used in this study to synthesize Bi-2223 is summarized in Fig. 6. Metal nitrates and acetates were used as the cationic sources and citric acid and ethylene glycol as the monomers for forming the polymeric matrix. Metal nitrates and acetates of the appropriate composition ratio were dissolved in a mixture of citric acid and ethylene glycol (1:8 molar ratio). A clear solution was then produced by heating the mixture at 90
C for 20 min. The resulting solution was further heated at 140
C to induce esterification and distil out excess ethylene glycol, which resulted in a substantial increase in solution viscosity. The viscous solution was then vacuum produced by calcining the precursors in air at various temperatures (300  800
C) for a few hours. In the work, the authors stated that the Pechini process is an attractive method to synthesize Bi(2223) superconductor materials for the high-Jc superconducting tape (Ag/Bi(2223)) and requires a much lower calcination temperature and shorter calcination time than the solid-state reaction. Polycrystalline Bi (2223) powder is composed of uniformly sized, ultrafine particulate with an average particle size of about 20 nm. Superconductivity investigation showed that the Pechini-synthesized materials display excellent performance.

Lithium Transition Metal Phosphates as Cathode Materials for HighVoltage Li-Ion Batteries Another example of application of the Pechini method concerns the preparation of Li-ion containing metal phosphates (LiMPO4, where M is a divalent transition metal, typically Fe, Co, Ni, Mn, and mixtures thereof) as cathode materials for (high-

1080 Fig. 5 Flowchart of the modified Pechini method to prepare high purity superconducting powder (Reproduced from Chiang et al. (1993))

L. Dimesso

Dissolve metal nitrate in water solution (except for thallium)

Add oxalic acid equal to 1/2 (total molar nitrate)

Titrate 25% ammonium solution to adjust pH = 6.5 – 7.5

Bismuth cation not fully complexed after stirring because of low solubility

Add Tl2O3 powder (D, E only)

Gel formation by solvent evaporation at 120°C

Organic decomposition at 250°C/2 hrs and 300°C/1 hr

Grinding of dried gel

Calcination at 750°C/12 hrs (A – C) of 700°C/3 hrs (D, E). Oxygen required for A, D, E.

Pellet-sintering. Oxygen required for A, D, E.

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Pechini Processes: An Alternate Approach of the Sol-Gel Method. . .

Fig. 6 Flowchart of the synthesis procedure (Reproduced from Peng et al. (1998) with permission of Springer)

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citric acid and ethylene glycol (1:8 molar ratio) ↓ dissolution 90 °C ↓ clear solution ↓ Bi(NO3)35H2O, Pb(AC)2 · 5H2O, Sr(NO3)2, Ca(AC)2 · H2O, Cu(NO3)2 · 3H2O ↓ dissolution 90 °C ↓ viscous solution ↓ estification 140 °C ↓ vacuum drying 180 °C ↓ estimation of ethylene glycol and polymerization ↓ polymeric precursor ↓ firing air, 300–800 °C ↓ fine powders

voltage) energy storage devices. Lithium-ion batteries have dominated the battery industry for the past several years in portable electronic devices due to their high volumetric and gravimetric energy densities. The success of these batteries in smallscale applications translates to large-scale applications, with an important impact in the future of the environment by improving energy efficiency and reduction of pollution. For more details on the progress that allows several lithium-intercalation compounds to become the active cathode element of a new generation of Li-ion batteries, namely, the 5-V cathodes, on the mechanisms which rule the performance of the batteries as well as the problems which are to be solved, I recommend to read the chapter in this book by Mosa (▶ Chap. 90, “Lithium Intercalation Materials for Battery Prepared by Sol-Gel Method”) and a review by Julien and Mauger (2013). I will summarize very shortly a few important concepts before to describe the advantage of the preparation of the Pechini (and Pechini-assisted) methods on the properties of the cathode materials. The most essential parameters in chemical energy storage devices (batteries) are specific energy, energy density (in both cases, the larger the better), cost (the lower

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the better), and safety. The cell-specific energy and energy density depend, first of all, on the cell chemistry, being reflected in its potential and charge capacity values. From this standpoint, Li-based cells hold much promise because Li metal is the most electropositive (E0 = 3.04 V vs. SHE) and light (ρ = 0.53 g cm3) material. However, employing Li metal in a secondary cell is challenging, since the possibility of dendrite growth poses risks of anode–cathode shorting. In the 1970s–1980s, the concept of a Li-ion cell (“rocking chair battery”) was demonstrated; this concept was based on the substitution of a Li metal anode with Li-ion intercalation compounds (Kraytsberg and Ein-Eli 2012). The lithium is in an “almost atomic” state in a carbonaceous anode material, and it is “almost Li+” state inside the cathode material, being oxidized by a transition metal redox couple. Whereas lithium mobility in the carbon anode is sufficiently high, the development of cathode materials with substantial Li+ mobility turned out to be an issue of prime importance. Various techniques have been used to prepare olivine structured LiMPO4 powders and composites including Pechini method and modified Pechini methods. For M = Fe our group (Dimesso et al. 2011a) reported the preparation of powders and composites using three-dimensional Vulcan carbon black prepared by a “the Pechini-assisted reverse polyol method” in which the LiFePO4 (LFPO4) powder was prepared by dissolving in water Li(CH3COO)2H2O (lithium acetate) and Fe (SO4)2*7H2O (iron(II) sulfate) as precursors with citric acid (2 x mol [Fe]), then adding DEG till the water weight percentage reached the desired value (10–40% wt). Finally, phosphoric acid in equimolar ratio with Li and Fe ions was added. The obtained solution was heated up to 105
C and kept at that temperature for 2–4 h. After cooling down the solution, the product was recovered by adding acetone and then separated by filtration. The precipitate was washed up by suspending it in acetone, then water, and acetone again and recovered by filtration more times. The cyclic voltammetric (CV) curves, shown in Fig. 7, of the pure-LFP do not reduplicate the processes occurring at the similar potentials. The peak voltage separations as well as the broadness of the CV curves indicate clearly that the electrochemical kinetics is strongly inhibited. On the other hand, in the LFP–VCB composites, the CV profiles reduplicate from the second cycle. The polarization in the first cycle is obvious and peak voltage separation is 0.23 V in the first cycle, while it is only 0.19 V from the second cycle. In the composite cathode, well-developed CV loop confirms that the kinetics of lithium intercalation and deintercalation are markedly improved compared to pure LiFePO4. A valid alternative to the carbon powder is the use of commercially available carbon foams. Indeed, they ensure good inter-particle conductivity and the continuous macroporous network allows an efficient transport route for the solvated ions. Our group extended the investigation to LiFePO4/carbon foam composites (Dimesso et al. 2011a, b, c). During this investigation, pure water, due to the rheological properties, favored the infiltration of the ions containing solution into the porous architecture of the foams. The starting solution, with a concentration of the precursors of 0.1 M, was heated up to 80–90
C and kept at that temperature for 2–4 h.

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Fig. 7 Cyclic voltammograms recorded for (a) LiFePO4, (b) LiFePO4 – Vulcan carbon black composites (Reproduced from Dimesso et al. (2011a) with permission of Springer)

After cooling down the solution, the product was separated by filtration. The commercial foams have been cut as disks having 5 mm diameter and 1–2 mm thickness. The composites were prepared by soaking the commercial foams in the starting aqueous solution at 70
C for 2 h. To obtain the olivine structured LiFePO4 phase, the powder was annealed at different temperatures ranging from 600
C to 700
C for 15 min. The micrographs in Fig. 8a show the morphological surface of the used foams. The porous architecture of the foams with hierarchical pore size distribution in micro-, meso-, and macropore ranges can be clearly recognized. We emphasize that in this “sponge” approach, the electrolyte layer is formed around a random 3-D network of electrode material. This design strategy also represents a concentric configuration in that the electrolyte envelops the electrode material, while the other electrode material fills the macroporous and mesoporous spaces. Short transport-path characteristics between the insertion electrodes are preserved with this arrangement. The micrographs of the CF–LFP composites prepared by soaking before and after annealing at 600
C for 10 min under nitrogen are shown in Fig. 8b, c. The

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Fig. 8 SEM pictures of (a) carbon foam as delivered; (b) carbon foam – LiFePO4 composites before annealing at T = 600
C; (c) carbon foam – LiFePO4 composites after annealing at 600
C for 10 min (Reproduced from Dimesso et al. (2011b) with permission of Springer)

micrograph shows a very homogeneous coating of the foam surface and consequently a more homogeneous morphology. By soaking, the foam surface is covered by a continuous layer of liquid in which the Li+, Fe2+, and (PO4)3 ions are uniformly distributed. The slow evaporation of the solvent leads to a “uniform”

Pechini Processes: An Alternate Approach of the Sol-Gel Method. . .

Fig. 9 (a) Discharge capacity, (b) discharge profile at C/25 discharge rate for carbon foam – LiFePO4 composite samples that were annealed at increasing temperatures (Reproduced from Dimesso et al. (2011b) with permission of Springer)

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layer on the foam surface (Fig. 8b). After annealing under nitrogen, the formation of a uniform layer of crystalline LiFePO4 can be observed (Fig. 8c). The discharge capacities for the carbon foam/LiFePO4 composites are presented in Fig. 9a. The annealing temperature had a significant effect on the capacity of the battery. The optimal temperature was determined to be 600
C with a capacity of 85 mAhg1 attained at a discharge rate of C/25. Although the sample annealed at 650
C delivered at the beginning a capacity of 110 mAhg1 at a rate of C/25, a capacity loss was observed during the cycling. This difference is even more evident at the slow discharge rate of C/10. This difference is likely due to the structural features of the composite materials. The discharge profiles in Fig. 6b have noticeably curved profiles even at such a low discharge rate of C/25. The 3.4 V voltage drops as the cell discharges due to polarization. Due to the promising results, we have extended the Pechini-assisted sol-gel process to other systems such as those containing Co and Ni (Dimesso et al. 2012a, b). We have reported the preparation of LNP and LCP by a

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Pechini-assisted sol-gel process that provides material exhibiting redox peaks at 5.2 and 4.9 V versus Li+/Li. Mg-substituted LNP/graphitic carbon foam composite was also synthesized by the same method, showing a discharge capacity of 126 mAhg1 at C/10 rate by substituting 0.2 Mg for Ni (Dimesso et al. 2013).

Concluding Remarks A short historical background, the basic principles, and the use of Pechini-type solgel methods for the synthesis of ceramic high-Tc superconductors and Li-ion containing olivine phosphates, including composites based on carbon nano-powder and foams, have been presented. Most of the difficulties and disadvantages that are often present in the alkoxide-based sol-gel process, such as high cost, unavailability, toxicity, and fast hydrolysis rate (thus difficult in controlling the homogeneity of different components during experimental processes) of the alkoxide precursors, can be avoided in the Pechini-type sol-gel process. Although in this chapter, I only presented the preparation of high-Tc ceramic superconductors and olivine structured phosphates as possible cathode materials for Li-ion batteries via the Pechini-type process, in view of its extensive applicability, many other kinds of optical and other functional materials with different forms by this method instead of the alkoxides sol-gel process can be prepared. Among these materials, films as planar waveguides (such as PbTi4O9), films with optical absorption and coloring (such as GeO2–V2O5 and TiO2–CeO2 systems), reflecting coating films (such as In2O3–SnO2, PbO–TiO2, and Bi2O3–TiO2), nonlinear optical films (LiNbO3, PbTiO3, and KTiOPO4), electrochromic films (such as WO3, V2O5, and V2O5–TiO2), and surface modification of porous and nonporous materials (such as SiO2, Al2O3, TiO2, etc.) with desired optical, electric, and magnetic properties for certain application purposes are included. In one word, the Pechini-type sol-gel process seems to be versatile and practical in the preparation of structural and functional materials with different forms. When the alkoxides sol-gel and other related processes cannot work well in some cases, one may get help from the Pechini-type sol-gel process to a great degree and might find some surprises. Acknowledgments I owe personally many thanks to Prof. W. Jaegermann (Head of the Surface Science Division, Earth and Material Sciences Department, Technische Universitaet Darmstadt) for giving me the possibility to contribute to this project and for the support during my work.

References Anderson HU, Pennell MJ, Guha JP. Polymeric synthesis of lead magnesiumniobate powders. Adv Ceram. 1987;21:91–8. Bednorz JG, M€uller KA. Possible high TC superconductivity in the Ba-La-Cu-O system. Zeitschrift f€ur Physik B. 1986;64:189–93.

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Chiang C, Huang YT, Shei CY. A summary on the modified Pechini method to prepare high purity ceramic superconducting powders. Chinese J Mater Sci. 1993;25:50–9. Cadogan IG, MacWilliam IC, Parson R et al. Dictionary of organic compounds, Vol. 2. New York: Oxford University Press; 1965. p. 715. Dimesso L, Spanheimer C, Jacke S, Jaegermann W. Synthesis and characterization of LiFePO4/3dimensional carbon nanostructure composites as possible cathode materials for Li-ion batteries. Ionics. 2011a;17:429–35. Dimesso L, Spanheimer C, Jacke S, Jaegermann W. Synthesis and characterization of threedimensional carbon foams–LiFePO4 Composites. J Power Sources. 2011b;196:6729–34. Dimesso L, Jacke S, Spanheimer C, Jaegermann W. Investigation on 3-dimensional carbon foams/ LiFePO4 composites as function of the annealing time under inert atmosphere J. Alloys Compd. 2011c;509:3777–82. Dimesso L, Jacke S, Spanheimer C, Jaegermann W. Investigation on LiCoPO4 powders as cathode materials annealed under different atmospheres. J Solid State Electrochem. 2012a;16:3911–9. Dimesso L, Becker D, Spanheimer C, Jaegermann W. Investigation of graphitic carbon foams/ LiNiPO4 composites. J Solid State Electrochem. 2012b;16:3791–8. Dimesso L, Spanheimer C, Jaegermann W. Effect of the Mg-substitution on the graphitic carbon foams – LiNi1yMgyPO4 composites as possible cathodes materials for 5 V applications. Mater Res Bull. 2013;48:559–65. Eror NG, Anderson HU. Polymeric precursor synthesis of ceramic materials. MRS Proc. 1986;73:571. https://doi.org/10.1557/PROC-73-571. Julien CM, Mauger A. Review of 5-V electrodes for Li-ion batteries: status and trends. Ionics. 2013;19:951–88. Kakihana M. Sol–gel preparation of high temperature superconducting oxides. J Sol-Gel Sci Technol. 1996;5:7–55. Kakihana M, Yoshimura M. Synthesis and characteristics of complex multicomponent oxides prepared by polymer complex method. Bull Chem Soc Jpn. 1999;72:1427–43. Kraytsberg A, Ein-Eli Y. Higher, stronger, better. A review of 5 volt cathode materials for advanced lithium-ion batteries. Adv Energy Mater. 2012;2:922–39. Lee H, Hong M, Bae S, Lee H, Park E, Kim K. A novel approach to preparing nano-size Co3O4-coated Ni powder by the Pechini method for MCFC cathodes. J Mater Chem. 2003;13:2626–32. Lessing PA. Mixed-cation oxide powders via polymeric precursors. Ceram Bull. 1989;68:1002–7. Lin J, Yu M, Lin C, Liu X. Multiform oxide optical materials via the versatile Pechini-Type sol–gel process: synthesis and characteristics. J Phys Chem C. 2007;111:5835–45. Massachusetts College of Fine Arts (MCLA) library. “The Sprague Log” Preserving a Company Newsletter, Hardman Library Grant Project, Massachusetts College of Liberal Arts, (1988–1989). http://mcla.libguides.com/localhistory/spraguelog Pechini MP. Patent US 3 330 697 (1967a). Pechini MP. Patent CA 759514 (1967b). Peng ZS, Hua ZQ, Li YN, Di J, Ma J, Chu YM, Zhen WN, Yang YL, Wang HJ, Zhao ZX. Synthesis and properties of the bi-based superconducting powder prepared by the Pechini process. J Supercond. 1998;11(6):749–54. Reichenbach HM, An H, McGinn PJ. Combinatorial synthesis and characterization of mixed metal oxides for soot combustion. Appl Catal Environ. 2003;44:347–54. Roy R, Osborn EF. Metal organics as ceramic precursors. Am Mineral. 1954;39:853–86. Roy DM, Roy R. Synthesis and stability of minerals in the system MgO-AI2O3-SiO2-H2O. Am Mineral. 1955;40:147–78. Tai LW, Lessing PA. Modified resin-intermediate processing of perovskite powders: part I. Optimization of polymeric precursors. J Mater Res. 1992;7:502–10. Tuttle BA, Voigt JA. In: Messing GL, Fuller Jr ER, Hausner H, editors. Ceramic powder science II. Westerville: American Ceramic Society; 1988. Chapter 4: Consolidation of ceramic thick films, pp. 62–69.

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Yoshino M, Kakihana M, Cho WS, Kato H, Kuto A. Polymerizable complex synthesis of pure Sr2NbxTa2-xO7 solid solutions with high photocatalytic activities for water decomposition into H2 and O2. Chem Mater. 2002;14:3369–76. Weber IT, Rousseau A, Guilloux-Viry M, Bouquet V, Perrin A. Microstructure comparison between KNbO3 thin films grown by polymeric precursors and PLD methods. Sol State Sci. 2005;7:1317–23. Zhang SC, Messing GL, Huebner W, Coleman MM. Synthesis of YBa2Cu3O7-x fibers from an organic acid solution. J Mater Res. 1990;5:1806–12.

Modified Pechini Synthesis of Oxide Powders and Thin Films

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Tor Olav Løveng Sunde, Tor Grande, and Mari-Ann Einarsrud

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . What is the Modified Pechini Synthesis? . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Step 1: Aqueous Solution of Precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Step 2: Drying and Heating of the Solution Obtained and Formation of a Polymeric Resin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Step 3: Decomposition of the Precursor Material to Obtain an Oxide Powder . . . . . . . . . . . Powders . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bismuth Ferrite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Indium Tin Oxide . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Phosphors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lanthanum Transition Metal Perovskite Oxides for Energy Technology . . . . . . . . . . . . . . . . . High-Temperature Superconductors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Cathode Materials for Li Ion Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thin Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Superconducting Thin Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Transparent Conducting Oxides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thin Film Phosphors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

The modified Pechini method has become one of the most popular synthesis methods for complex oxide materials due to its simplicity and versatility. The method can be applied to synthesize nanocrystalline powders, bulk materials, as T. O. L. Sunde (*) Department of Sustainable Energy Technology, SINTEF Materials and Chemistry, Oslo, Norway e-mail: [email protected] T. Grande · M.-A. Einarsrud Department of Materials Science and Engineering, NTNU Norwegian University of Science and Technology, Trondheim, Norway e-mail: [email protected]; [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_130

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well as oxide thin films. Here, we present a comprehensive review of the method with focus on the chemistry through the three stages of the process: preparation of stable aqueous solution, polyesterification to form a solid polymeric resin, and finally decomposition/combustion of the resin to form an amorphous oxide followed by crystallization of the desired oxide phase. The review include several examples of important technical oxide materials where the method has been successfully been applied to prepare oxide powders and bulk or thin films.

Introduction Oxide materials in form of particles and bulk or thin films are continuously becoming more important and thereby have also synthesis methods of oxides grown to be an important research field including solution-based routes. Pechini described in a patent from 1967 (Pechini 1967) how films of titanate and niobate dielectrics could be prepared by utilizing the ability of certain alpha-hydroxycarboxylic acids, such as citric, lactic, and glycolic acids, to form polybasic acid chelates with different cations. The chelates underwent polyesterification when heated in a polyhydroxy alcohol solution, resulting in a transparent solid resin, which maintain the desired homogeneity of the cation distribution from the solution. Upon calcination of the resin, the organics are removed leaving the desired oxide ceramic composition as the residue. The precursors Pechini suggested were oxides, hydroxides, alkoxides, and carbonates. The synthesis method first described in the Pechini patent has further been modified by several authors. Marcilly et al. (1967, 1970; Courty et al. 1973) proposed the same year as Pechini a similar approach where the polyol (ethylene glycol) was completely replaced with water forming an amorphous gel-like matter instead of the polymer. This process is termed amorphous citrate or metal complex method. An advantage with this process was the lower amount of organics, but the homogeneity of the resulting oxide seemed to be lower than the original Pechini route. The following development lead to the use of other precursors, e.g., nitrates containing crystallization water, like in the synthesis of LaMnO3+δ using the polymerizable complex route and nitrate salt precursors (Kakihana et al. 1999). Focus on reducing the amounts of organics occurred in the following years, and water was introduced as a solvent for the precursors. One of the first publications using the water-based approach was a study of superconducting film formation (Chiang et al. 1991). The term “modified Pechini process” is today mostly used for a process where aqueous solutions are used as precursors and this can be defined in the following way: The modified Pechini type of synthesis process to oxide materials starts with a homogeneous aqueous solution containing the desired cation precursors in stoichiometric ratio and selected additives, which by evaporation and reactions is converted to a rigid cross-linked polymer hindering segregation of the cations. The polymer is further converted to a homogeneous oxide powder or film by heat treatment. The modified Pechini synthesis method is simple and does not need sophisticated laboratory infrastructure facilities. Aqueous processing routes to oxides are also attractive related to environmental concerns. The synthesis can be performed in a

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simple beaker using a hot plate, and the heat treatment can be done in a regular laboratory furnace. The modified Pechini process is hence a robust and versatile synthesis method used for many different applications. Due to the fact that the homogeneity from the aqueous solution is maintained through the polymeric resin into the final oxide material, the method is well suited for the preparation of the wide range of homogeneous multicomponent oxide materials. The method can easily be extended to chemical solution deposition of thin films or coatings. There are in principle no restrictions to the number of cations in the solutions, and this simple method actually becomes more attractive the more complex the product is. Here we provide a detailed review of the chemistry related to the modified Pechini process to oxide powders and films. We focus on the three different steps of the synthesis process including solutions to challenges that might appear using the method. Finally, the applicability of the process is illustrated by describing several selected examples of important technical oxides where the synthesis is applied with success.

What is the Modified Pechini Synthesis? The three major steps of the modified Pechini process to achieve a homogeneous oxide powder or film are illustrated in Fig. 1. In the following, a thorough description of the chemistry related to these three steps is provided.

Step 1: Aqueous Solution of Precursors A stable aqueous chelated solution with the cation precursors is paramount for a successful wet chemical synthesis. The complexing agent polybasic hydroxy carboxylic acid forms the polybasic carboxylic acid chelates. The solution needs to be stable over time preventing precipitation, as this will introduce inhomogeneities in the resulting material. The cation precursors should be soluble in water or in aqueous solutions with the chelating agents; hence, typical precursors can be hydroxides, alkoxides, acetates, chlorides, citrates, and nitrates. The most economically abundant and highly soluble precursor for most metals is their nitrates. Concentration of the cations will normally be in the range from 0.1 to 1.0 M. It is of imperative importance that the cations are mixed in the correct stoichiometric ratio. Thermogravimetric analysis or ICP is normally performed on each stock solution of cation precursors to determine and control the exact concentration. The most common polybasic hydroxyl carboxylic acid used is citric acid. The molecular structures of citric acid together with some other relevant polybasic hydroxyl carboxylic acids are given in Fig. 2. The middle carboxylic acid group of citric acid is the most acidic due to the electron withdrawing power of the alpha-OH group, and hence, this carboxylic acid group will form the strongest complexes. The degree of protonation of citric acid is dependent on the pH with the three pKa values being 3.13, 4.76, and 6.39.

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Fig. 1 Flowsheet showing the different steps of the modified Pechini process producing a complex oxide from an aqueous solution of precursors and additives

Fig. 2 Structure of relevant polybasic hydroxyl carboxylic acids frequently used during the modified Pechini synthesis

Most transition metals form stable complexes with the citric acid and other similar chelating agents which is common for all acidic cations. However, basic cations like the alkaline earths form weak complexes. The stability of the complexes is dependent on pH and concentration, and the stability of several important cations can be

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found from data provided by standard databases (The IUPAC Stability Constants Database, Academic Software). Basic cations are found in the lower left corner of the periodic table, while the more acidic ones are found in the right upper part of the periodic table. The cation charge, cation size, as well as number of valence electrons are important factors determining the acidity or basicity of cations and hence the stability of chelate complexes. For the most basic cations like Ba, it has been found necessary to add EDTA as an additional complexing agent (Sletnes et al. 2016). Normally the carboxylic acid like citric acid is added in excess, typically with a molar ratio citric acid:cations from 1 to 3, and this ratio is important for the success of the synthesis as will be demonstrated in some of the examples provided (Fontaine et al. 2004; Nityanand et al. 2011). The order of mixing of the cations and chelating agents has been studied in a few examples, and it is normal to add the cation forming the most stable chelate first followed by the cations forming less stable chelates (Sletnes et al. 2016). The structure of the chelate complexes has been revealed in some cases, but this is not a topic that has gained much focus in the literature. For a few materials, like BaTiO3 and Y2Ti2O7, it has actually been indicated that heterometallic chelate complexes can be formed, where both cations are bound with a stoichiometric ratio in the same complex (Arima et al. 1996; Kakihana et al. 1996). This gives a very high degree of homogeneity, which is beneficial for the preparation of these compounds. The size of the polybasic acid and the ratio between the acid and cations have shown to be of importance as more organics has to be burned off, influencing on the local temperature as well as the partial pressure of oxygen (Selbach et al. 2007). In a synthesis protocol closely related to the modified Pechini synthesis, a polymer (PVA or PEG) is used to physically stabilize the cations by a proposed entanglement around the cations. By this polymer-complex solution method, less amount of organics is used, and for the synthesis of CaAl2O4 a phase-pure material was prepared with a higher amount of cations than the number of functional groups in the polymer (G€ulg€un et al. 1999). For some specific cations, it has more recently been observed that complexing agents with –OH groups like ethylene glycol and polyvinyl alcohol (PVA) are better compounds to promote homogeneity of the final oxide than carboxylic acids. This has empirically been found to be of great advantage for amphoteric cations like In3+, Sn4+, and Bi3+ in the preparation of indium tin oxide (Kundu and Biswas 2008; Sunde et al. 2012) and BiFeO3 (Selbach et al. 2007; Liu et al. 2010). An amino carboxylic acid, e.g., glycine, has also been used as a complexing agent and fuel in a related synthesis route, glycine nitrate method (Chick et al. 1990). To promote the polymerization into a resin, a polyalcohol is added, normally ethylene glycol, to promote polymerization with the polybasic carboxylic group giving a polyester, according the chemical reaction given in Fig. 3. The amount of ethylene glycol added should at least be the minimum amount necessary with respect to the esterification reaction. The role of the polyalcohol is questioned in the literature and historically a large excess was used (Pechini 1967; Kakihana et al. 1996), but now it is common to add polyalcohol in the same molar ratio as the carboxylic acid. The amount of ethylene glycol has been shown to be important

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Fig. 3 Illustration of the chemical reaction between a polybasic carboxylic acid chelate and ethylene glycol during the formation of a complex perovskite oxide (Kakihana and Yoshimura 1999). In the modified Pechini process, water has to be evaporated in order to form the polymeric resin. Figure reprinted with permission from Bull. Chem. Soc. Japan

for the morphology of ZnO powder produced by the method (Farbun et al. 2013). The polyesterification reaction is dependent on the presence of protonated carboxylic groups and hence the pH of the solution (Tai and Lessing 1992a; Kakihana and Yoshimura 1999). Sometimes an acid (typically nitric acid) has to be included in the solution to catalyze the polymerization dependent on the original pH. On the other

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hand, the addition of a base (typically ammonium hydroxide) to increase the pH can sometimes be necessary to avoid precipitations of for instance barium nitrate (Chu and Dunn 1987).

Step 2: Drying and Heating of the Solution Obtained and Formation of a Polymeric Resin The stable aqueous solution of the cation precursors and the chelating agents prepared in step 1 is further dried at temperatures in the range 110–150  C to remove water. As the amount of water decreases, it is important to completely prevent precipitation of salts as this will be detrimental to the homogeneity of the final material. The viscosity is increasing during the evaporation of water, and the solution is finally turned into a polymeric-like resin without any precipitates. Basic cations might precipitate as nitrates, chlorides, acetates, as well as citrates from the chelating agent during evaporation of water. Possible precipitation reactions should be carefully considered during the initial design of the synthesis. Upon further heating of the polymeric resin, a large increase in volume (foaming) occurs due to release of water from hydrated chelates, decomposition of complex anions from the precursors, as well as decomposition of the organics. The foaming reduces segregation of cations during the process due to the formation of long diffusion paths in the very fine foam formed by the gas development. After drying and reactions, the aqueous precursor solution is turned into a brittle voluminous spongy-like precursor material.

Step 3: Decomposition of the Precursor Material to Obtain an Oxide Powder The third step is thermal decomposition of the precursor material to remove the organic part of the resin. Thermal decomposition of the resin usually occurs below 400  C. The decomposition products are normally not identified and not well described in relevant literature. The spongy-like precursor material is first turned into an amorphous material that is further crystallized into a single-phase target material (Tai and Lessing 1992a; Sletnes et al. 2016). The crystallization temperature is dependent on the system but can be as low as 400  C as, for example, in ITO (Sunde et al. 2012). The decomposition is exothermic, sometimes causing the decomposition and crystallization to occur simultaneously, thereby making it challenging to identify the real crystallization temperature (Sunde et al. 2012). Metastable oxide phases may also crystallize from the homogeneous amorphous oxide due to the low crystallization temperature of the precursor. This is an interesting aspect with this synthesis approach. Upon further heating, the metastable phase will be transformed to the thermodynamically stable phase (Schumm et al. 2011; Sunde et al. 2012). During preparation of oxides with very basic cations, the formation of intermediate carbonates might occur due to the availability of CO2

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from the decomposition of the organics. These carbonates will decompose upon further heating forming the target phase. Several authors also claim that carbonates are not formed during synthesis of, for example, BaTiO3 and SrTiO3 and that the homogeneity is maintained throughout the synthesis (Cho et al. 1990; Arima et al. 1996; Kakihana and Yoshimura 1999).

Film Deposition The aqueous chelated solution developed under step 1 can also be utilized for deposition of thin films by the modified Pechini process. The film is then deposited on a substrate by spin coating or dip coating, and the necessary equipment is significantly simpler and less expensive than for physical deposition techniques. This precursor film is dried on a hot plate, pyrolyzed (heat treated) to decompose organic additives and complex anions from the precursors, and crystallized. The two latter steps are often combined in a single heat treatment step, normally performed by rapid thermal processing. The microstructure of the resulting film is very dependent on the type of nucleation, being dependent on the pyrolysis temperature, heating rate, and crystallization temperature (Schwartz et al. 2004). The thickness of each deposited film can tailored from a few to a few hundred nanometers by varying the cation concentration and viscosity of the solution or by changing the spinning or dipping speed (Bernardi et al. 2002; Sunde et al. 2014). The total thickness of the film can also be increased by repeating the deposition procedure. Health and Safety Concerns The Pechini process is a low temperature, simple, and robust method for making oxide materials; however, there are some HSE issues that should be considered. Starting with precursors containing a strong oxidation agent like the nitrate ion could actually lead to ignition of the organic material during drying and heating giving a combustion reaction. Care should therefore be taken when using these precursors by testing a small batch of the synthesis. The drying and heat treatments in this case should be under controlled conditions. Poisonous gases might also evolve during the decomposition of the precursor anions, e.g., nitrous gases from the decomposition of nitrates, which should be used under proper ventilation. Moreover, during the drying of the solution, a large volume increase will occur due to foaming caused by gas evolution. It is important to use large enough container to take this into consideration. Using a polybasic amino carboxylic acid as complexing agents and fuel as in the glycine nitrate method might lead to a combustion dependent on the nitrate/ glycine ratio. Special care must be taken in doing this synthesis and only small amount can be made in each batch. Finally, it should also be emphasized that the use of water as a solvent in contrast to toxic and/or expensive organic solvents is beneficial with respect to environmental as well as economic concerns. Most of the precursors are also environmentally friendly. Solutions to Frequent Challenges Correct stoichiometry not obtained: The cation concentration and ratio in the solution has to be accurate, especially for line compounds with very limited solid

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solubility. Several cation precursors like nitrates contain uncertain amount of crystal water and calibration of the cation content in the solution has to be carried out. Solid precursors (alkali nitrates, oxides) also contain adsorbed water and needs to be dried before use. For cations forming volatile species, evaporation below the polymeric resin decomposition/burning may result in incorrect stoichiometry. Precipitation of crystalline phases before the formation of the polymeric resin: Make sure that complexation of the less-soluble salts is taken into account to prevent precipitates like nitrates, oxalates, citrates, acetates, or tartrates especially in case of basic cations. If this is the case, change the cation precursor or the chelating agent. Inhomogeneities in the oxide powder/film: Carbonates might form during the heat treatment especially of basic cations like Ba. These carbonate phases will disappear at higher temperature due to reaction with the residual oxides, but a single-phase product might be challenging to obtain due to the initial formation of several phases. In some cases, the formation of thermodynamically stable inert compound might occur, e.g., formation of parasitic pyrochlore formation instead of the desired perovskite phase (Rørvik et al. 2009). Redox reactions between the cations and the organics may also in some cases give elemental metal like Bi (Leonard et al. 2002), which may be prevented by changing in the type of cation precursors or complexing agents.

Powders Bismuth Ferrite Bismuth ferrite, BiFeO3, is a multiferroic material due to the coexistence of ferroelectricity and antiferromagnetism. The material has received considerable attention in the last decade due to potential applications in data storage, sensors, and devices for spintronics (Wang et al. 2003). BiFeO3 is known for being difficult to prepare by solid-state reactions due to its relatively low melting point (incongruently melting at 934  C) and preferential evaporation of Bi2O3 at high temperatures. Even small offsets in stoichiometry can lead to secondary phases of Bi25FeO40 and Bi2Fe4O9 (Selbach et al. 2009). A variety of wet chemical synthesis routes have been reported (Kim et al. 2005), but in most reports only nearly phase-pure materials were prepared (Shetty et al. 2002). Several modified Pechini-related synthesis routes have been used, with different polybasic carboxylic acids as complexing agents with and without the addition of EG as polymerization agent (Selbach et al. 2007). Here, phase-pure BiFeO3 was obtained by tartaric and malic acid, with and without EG, and maleic acid with EG. It was proposed that a requirement for the formation of phase-pure materials was the presence of both COOH groups for complexing Bi3+ and Fe3+ and OH groups for polyesterification. A literature review from Liu et al. (Liu et al. 2010) noted that it appeared that the presence of the OH-group is critical for the formation of phase-pure BiFeO3, while the presence of the COOH groups does not seem to be significant, as can be seen from Table 1. Synthesis routes with PVA and EG have

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Table 1 Literature review of modified Pechini syntheses where different complexing agents have been used for the preparation of BiFeO3. (Modified from Liu et al. 2010) PVA EG

# of OH X 2

# of COOH 0 0 2 2 3

Phase-purity ✓ ✓ Without EG ✓ ✓ 

Tartaric Malic Citric

2 1 1

Maleic Succinic Malonic Oxalic EDTA

0 0 0 0 0

Reference (Liu et al. 2010) (Park et al. 2007) With EG ✓ ✓ ✓

2 2 2 2 4

    ?

✓   ? 

(Ghosh et al. 2005a) (Selbach et al. 2007) (Jiang et al. 2006; Popa et al. 2007) (Selbach et al. 2007) (Selbach et al. 2007) (Selbach et al. 2007) (Ghosh et al. 2005b) (Liu et al. 2010)

✓: phase-pure; : not phase-pure; ?: not reported

been successful, while synthesis routes with maleic, succinic, malonic, and oxalic acids (acids with no OH groups) were not. Even the utilization of EDTA, known as a very strong complexing agent with four carboxylic acid groups, was unsuccessful (Liu et al. 2010). It is interesting to note that the classical Pechini synthesis route (Pechini 1967) with citric acid and ethylene glycol has been reported both to give phase-pure (Popa et al. 2007) materials and to give secondary phases (Ghosh et al. 2005a). However, in the successful experiment, EG was used as a solvent, i. e., in much larger quantities than when it is only added in smaller amounts as a polymerization agent to an aqueous solution. The majority of works use nitrates of bismuth and iron as cation precursors (Selbach et al. 2007). The choice of precursors and amount of organic additives can be important. Hardy et al. reported a Pechini-like synthesis where nitrates and citrates were used as precursor (Hardy et al. 2009). This combination can lead to a self-combustion of the gel during calcination, in combination with the evaporation of large amounts of decomposition gases, like CO (Hardy et al. 2005). This can locally create quite high temperatures and a reducing atmosphere, which again can lead to the formation of metallic bismuth. This phase segregation implies a loss of homogeneity during calcination which will be a disadvantage, even if the metal is readily oxidized in the later stages of calcination.

Indium Tin Oxide Tin-doped indium oxide, known as indium tin oxide (ITO), has been the most widely used transparent conducting oxide to date (Ginley and Perkins 2011). Due to excellent properties, ITO has found numerous technological applications such as flat panel displays, touch panels, energy-efficient windows, and solar cells. ITO thin films are industrially prepared by sputtering, where targets with high density are

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required (Gehman et al. 1992). It has been known for decades that it is difficult to obtain In2O3 and ITO with high density (Vojnovich and Bratton 1975; Nadaud et al. 1994), which has caused a considerable interest in the synthesis and sintering of nanocrystalline In2O3 and ITO. In most of these studies, nanocrystalline In2O3 and ITO were prepared by coprecipitation (Kim et al. 2002, 2006), although a few reports with modified Pechini methods exist (Sunde et al. 2012). The modified Pechini methods avoid the simultaneous precipitation of both indium and tin hydroxides, which ruins the homogeneity on the atomic scale using coprecipitation. Indium nitrate and tin chloride are the most commonly used precursors (Psuja et al. 2007), although alternatives like tin acetate have also been applied (Sunde et al. 2012). Abbas et al. prepared ITO powder from an aqueous solution with CA and EG (Abbas et al. 2014). CA was used with a molar ratio of 2:1 compared to the cations, while (EG) was added with a ratio of CA:EG of 80:20, 70:30, or 60:40. Powders prepared with the 80:20 ratio had the best sintering properties, possibly due to less agglomeration. Sunde et al. performed a series of experiments with a cation concentration of 0.5 M and different complexing agents, applied alone or in combination, namely, acetic acid, ethanol, ethylene glycol, succinic acid, and tartaric acid (Sunde et al. 2012). Here, it was observed that precipitations of organometallic salts occurred for almost all syntheses during the drying of the gel. Amorphous gels were only be obtained in the synthesis with EG and tartaric acid and EG in combination. This is an indication that the ROH-group is more efficient at chelating indium and tin cations than the RCOOH group. In another work, phase-pure ITO was prepared by a citrate-combustion method, without ROH groups (Wang et al. 2010), indicating that also the carboxylic acid group can also effectively chelate these cations when the experimental conditions are favorable. The organics in the obtained amorphous gels will decompose upon calcination, and the exothermic nature of the reaction causes the decomposition and crystallization to occur simultaneously already at 400  C (Wang et al. 2010). In some cases, also the metastable rhombohedral polymorph of ITO can be obtained during calcination at these low temperatures. This has been observed both when smaller organic complexing agents were used, giving less reducing atmospheres during calcination (Sunde et al. 2012), and when ZrO2 was used as a dopant (Abbas et al. 2014). An indication of the excellent homogeneity provided by the modified Pechini method is that phase-pure ITO powders can be prepared with significantly higher amounts of tin doping than the thermodynamic solubility limit (Sunde et al. 2013). The phase purity will remain until the temperature is high enough to activate cation mobility to give phase segregation. This can be seen in Fig. 4, where the material is phase-pure up to 800  C and a secondary phase of SnO2 appears at temperatures above 1000  C.

Phosphors Phosphors are materials which can emit light, and one of their emerging applications is as white light-emitting diodes (WLED), with the potential to significantly reduce

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Fig. 4 XRD of ITO powders with 100 cation% tin doping prepared from a modified Pechini synthesis after heat treatment at different temperatures. The material is phase pure at low temperatures, but an exsolution of SnO2 is observed from 1000  C and upwards (diffraction lines from SnO2 marked by *). (From Sunde et al. 2013)

the energy consumption of illumination (Shur and Žukauskas 2005). The phosphors often consist of host materials with complex composition doped with small amounts of the active light-emitting species and are sensitive to phase-purity, crystallinity, and impurities (Smet et al. 2011). Hence, wet chemical synthesis methods are highly applicable for their synthesis (Gai et al. 2014). The literature on phosphors prepared by modified Pechini methods is extensive, including host materials such as CaIn2O4 (Liu et al. 2007), Ca8La2(PO4)6O2 (Shang et al. 2012), YVO4 (Serra et al. 2000), YNbO4 (L€ u et al. 2015), and SrZrO3 (Zhang et al. 2008). In some of these reports, the materials are also prepared by solid-state reactions for comparison, and the modified Pechini-powders are shown to have equal or better properties (Wang et al. 2008b). A much studied group of novel red phosphor materials for WLED are Eu3+-doped oxides of Mo and W, including compounds such as NaLa(WO4)(MoO4) (Li et al. 2013), La2(WO4)3 (Kodaira et al. 2003), and double perovskites within the (Sr,Ba)2Ca(W,Mo)O6 series (Ye et al. 2011). The most common synthesis method of these materials is solid-state reaction, but there are several reports of modified Pechini synthesis being used (Zalga et al. 2011). However, it is not trivial to obtain phase-pure materials and parameters such as choice of complexing agent, pH, and calcination temperature must be optimized. In most of the reports, CA and EG are used as complexing and polymerization agents (Kodaira et al. 2003; Wang et al. 2008b), but other complexing agents have also been used. In a work by Sletnes et al., NaLa(WO4)(MoO4) was synthesized by using malic and tartaric acid, with and without EG (Sletnes et al. 2016). All of the syntheses yielded phase-pure materials after calcination at 600  C, but when malic acid was used alone, precipitation of a salt was observed in the gel. Although crystallization occurred already at 400  C, 600  C was necessary to completely oxidize and remove the organic residue. Increasing the calcination temperature increased the crystallinity and particle size, which often improves the optical

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properties. Interestingly, the best optical properties of the powder were obtained after calcination at 600  C, possibly due to some finite size effects. The same work also showed how phase-pure Ba2CaMoO6 could be synthesized by the utilization of CA and EDTA as complexing agent. Preliminary experiments with MA, TA, CA, EDTA, and DTPA, with or without EG, demonstrated the challenges related to the synthesis of this material, mostly due to precipitation of Ba(NO3)2 in the gel. Precipitations were observed when MA, TA, and EDTA were used, which further lead to secondary phases of Ba1-xCaxCO3 and BaMoO4 after calcination. Amorphous gels, indicating a good homogeneity, were obtained by using CA, DTPA, or EDTA and CA in combination. Among these, phase-pure materials were only obtained by the latter two routes at calcination temperatures higher than 800  C. Of the two routes, the one with EDTA and CA was deemed most promising, due to the suspected toxicity of DTPA. There are several modified Pechini syntheses of perovskites reported in the literature in which EDTA and CA are used in combination, and many of them involve basic cations, such as Sr, Ba, and La (Ding et al. 2008; Patra et al. 2011). The application of these two complexing agents in combination is clearly favorable, but the actual mechanism is not understood. It might be related to their favorable complexing, but also that the combination is a good combustion reagent, which increases the local temperature and the reaction rate. The difference in the required complexing agents for the synthesis of the two materials can be partly explained by the solubility of the precursors and the stability of the cation complexes. The stability constants for complexes with Ca2+, Ba2+, Na+, and La3+ are plotted in Fig. 5. The stability constants are significantly higher for EDTA and DTPA than for MA, TA, and CA, and the stability increases in the order Na+ < Ba2+ < Ca2+ < La3+. For NaNO3, the solubility is so high that strong complexation might not be necessary to achieve phase-pure materials. For Ba (NO3)2, on the other hand, strong complexing agents like EDTA should be used.

Lanthanum Transition Metal Perovskite Oxides for Energy Technology Several transition metal perovskites with lanthanum/strontium on the A-site have been extensively investigated as electrodes for solid oxide fuel cells (SOFC) (Tao and Irvine 2004; Baumann et al. 2006; Jiang 2008), such as La1-xSrxCoO3 (LSC), La1-xSrxMnO3 (LSM), La1-xSrxCr0.5Mn0.5O3 δ (LSCM), and La1-xSrxCo1-yFeyO3 δ (LSCF). Doped LaCrO3 has also received considerable attention as an interconnect material for SOFCs (Zhu and Deevi 2003). Tailoring the stoichiometry, microstructure, and phase purity of the powders is vital in order to optimize their performance, and powders of high quality is an essential prerequisite for these applications. The typical cation precursors for these perovskites are nitrates (Magnone et al. 2007), occasionally prepared by dissolving the parent oxide in HNO3 (da Conceição et al. 2009), although sometimes carbonates are used for the more basic cations (Fan and Liu 2009). Gupta and Whang also did a series of experiments

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Fig. 5 Stability constants (log K) of complexes of Ca2+, Ba2+, Na+, and La3+ with malic acid, tartaric acid, citric acid, EDTA, and DTPA. (From Sletnes et al. 2016)

where they changed between nitrates and acetates as precursors for the preparation of La0.9Sr0.1Cr0.85Fe0.05Co0.05Ni0.05O3 δ (Gupta and Whang 2007). They observed that using acetates gave less secondary phases, and discussed this in relation to the pH in the solution and its influence on the chelating ability of CA. Several complexing agents have been used, but the most typical is CA, both with (Gaudon et al. 2002) and without (Fan and Liu 2009) addition of EG as a polymerization agent. The ratio between CA and metal cations has been demonstrated to be an important parameter. For ratios less than three, precipitation of unidentified phases was observed, both for LSM (Da Conceição et al. 2011) and LSCF (Nityanand et al. 2011). On the other hand, by adding ammonia to adjust the pH to 8, Fan et al. prepared phase-pure LSCF with a CA:cation ratio of only 1.5 (Fan and Liu 2009). Strontium carbonate has been observed as a secondary phase for calcination at temperatures of 500–600  C due to the high stability of SrCO3, but phase-pure materials are obtained by increasing the calcination temperatures (Shao et al. 2009). This is illustrated by the XRD diffractograms in Fig. 6. It is interesting to note that lanthanum manganite typically can be made phase pure at lower temperatures than lanthanum strontium manganite, because the possible formation of SrCO3 can be avoided (Kakihana et al. 1999). The necessity of adding EG as a polymerization agent is also not clear. Phase-pure LSCF has been prepared by the amorphous citrate synthesis, without the use of EG (Magnone et al. 2007; Fan and Liu 2009). On the other hand, Kakihana et al. demonstrated that the polymerization was crucial in order to avoid secondary phases during the synthesis of lanthanum manganite (Kakihana et al. 1999). It appears that the polymerization is not essential, but it improves the homogeneity during the synthesis, thereby enabling the formation of phase-pure

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Fig. 6 XRD diffractograms of LSCF prepared with CA and EDTA as complexing agents after calcination at different temperatures. (From Shao et al. 2009)

materials at lower temperatures. This can be important if a high surface area of the powder is desirable, for instance, to provide a large surface area of a cathode material or enhance the sintering properties. Shao et al. performed an interesting series of experiments to shed light on the chelating of the precursors of LSCF (Shao et al. 2009). Here, CA, EDTA, or the combination of the two were used in experiments with all four cations of LSCF together and for the individual cations alone. Based on visual observations of precipitations and FTIR and XRD experiments, they observed that CA could form complexes with all four cations, but could not totally coordinate them (a molar ratio of 1.5 compared to the cations was used). On the other hand, EDTA could only form stable complexes with Co and Fe, not with La and Sr (molar ratio of 0.5). The best results were obtained by using both complexing agents together. The infrared spectrum of the LSCF complex precursor with CA and EDTA was not a pure superposition of the spectra of the individual cation complexes, indicating a strong interaction between the cations during the chelation. Similar observations have been reported in the synthesis of BaTiO3 (Kakihana and Yoshimura 1999) and La2Mo2O9 (Rocha and Muccillo 2003) using modified Pechini syntheses. In the early 1990s, Tai and Lessing performed a thorough fundamental investigation of the synthesis of La0.85Sr0.15CrO3 by the modified Pechini method (Tai and Lessing 1992a, b) where the ratio between CA and EG, the ratio between the organic additives and cations, as well as the amount of excess water were investigated. The optimal ratio between CA and EG in order to get a polymerization is between 40% and 57% CA. Three drops of nitric acid was added for each 100 mL of water, in order

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to catalyze the esterification reaction. The CA and EG should be mixed together before adding the cation precursors. If not, they could observe precipitations of strontium nitrate. If the ratio between organics and cations is too high, > 4, fierce ignition of the gel occurred during calcination, thereby increasing the temperature and creating hard agglomerates. It is also desirable to keep the amount of added organics low from an economic perspective. However, if the ratio was too low, < 0.5, the complexation was not sufficient to avoid secondary phases. A total of 150 mL water was necessary to aid the dissolution of CA when 1 mol of CA and 1 mol of EG were used. Excess water beyond this had two drawbacks. First, the evaporation of water also removed some quantities of the polymeric substance out from the boiling solution. And second, the formation of energetic water vapor during the evaporation could break the polymeric network that was already forming in the solution.

High-Temperature Superconductors Superconductors are materials which exhibit zero electrical resistance and expulsion of magnetic fields below a critical temperature, Tc. The discovery of high-temperature superconductors in copper oxide-based material systems with a Tc above 77 K (can be cooled by liquid nitrogen) in the late 1980s created enormous excitement and these materials were intensively investigated (Bednorz and M€uller 1986). Particularly oxides in the Y-Ba-Cu system (YBCO), like YBa2Cu3O7-x, received a significant amount of interest (Wu et al. 1987). There are numerous reports on different synthesis routes to prepare YBCO powder, as reviewed by Kakihana and Pathak and Misra (Kakihana 1996; Pathak and Mishra 2005). The powder characteristics, like homogeneity, phase purity, and particle size, will have a large influence on the properties of the final device, and both solid-state (Cava et al. 1987) and wet chemical synthesis routes (Kamat et al. 1991) have been employed for the synthesis of YBCO. Different sol-gel syntheses to prepare YBCO have been applied, and they are argued to give better homogeneity and be less complicated than other wet chemical methods, such as coprecipitation (Kakihana et al. 1989). The cation precursors used in modified Pechini syntheses for YBCO can either be nitrates (Kakihana et al. 1991) or Y2O3 and CuO dissolved in HNO3 together with BaCO3 (Lee et al. 1989). A possible precipitation of Ba(NO3)2 could be avoided by adding ammonia to adjust the pH to about 6 (Chu and Dunn 1987). In all of the reports, CA and EG were employed as complexing and polymerization agents. The ratio between metal cations and CA varied from 1:0.33 (Kakihana et al. 1991) to 1:6.5 (Shiomi et al. 1993). After evaporation of the solvent, the gels were calcined and phase-pure YBCO was obtained at around 900  C (Mazaki et al. 1991). Sometimes BaCO3 could be detected after calcination at lower temperatures (Lee et al. 1989). The resistivity of YBCO prepared with CA and EG is shown in Fig. 7, demonstrating the superconducting transition. The narrow transition

Modified Pechini Synthesis of Oxide Powders and Thin Films

resistivity (arbitrary unit)

36

1105

92 K

90 K R=0 70

80

90

100 temperature

110

120

Fig. 7 Resistivity as a function of temperature for a polycrystalline sample of YBa2Cu3O7-δ prepared by a modified Pechini method with CA and EG. (From Kakihana et al. 1991)

temperature interval is an indication that YBCO with high purity and structural homogeneity could be prepared (Kakihana et al. 1991) by this method. In addition to the modified Pechini routes, there are also several reports on amorphous citrate-syntheses to prepare YBCO (Yang et al. 1989) without the use of EG. Also oxalic acid has been used as the complexing agent (Sanjines et al. 1988). In some of these reports, Y2O3, together with barium and copper carbonate, was dissolved directly in citric acid, without the use of HNO3 (Karen and Kjekshus 1994). The molar ratio between cations and CA was typically 1:2 (Sanjines et al. 1988). Also in these syntheses the pH was often increased to prevent precipitation of Ba(NO3)2. Liu et al. argued that it was advantageous to use ethylenediamine to adjust the pH instead of ammonia, as it could react with nitric acid and thereby further prevent precipitation of Ba(NO3)2 (Liu et al. 1989). Another reported way to mitigate the issue with Ba(NO3)2 precipitation is to employ a combination of CA and EDTA as complexing agents (Van der Biest et al. 1991). Here, one solution is prepared where Y3+ and Cu2+ are complexed by CA and another with Ba2+ complexed by EDTA. After mixing the two solutions and adjusting pH to 7, the precipitation could be avoided (Niou et al. 1992).

Cathode Materials for Li Ion Batteries Lithium ion batteries have become increasingly important and have been at the focus of intense research the last decades. They are today routinely used in portable electronic devices and also show great promise for storage of energy from renewable sources and their use in electronic vehicles (Bruce et al. 2008; Goodenough and Kim 2010). The anode is typically graphite, whereas the cathode traditionally was LiCoO2 (Mizushima et al. 1980). Later a variety of potential cathode materials has received interest, among them several materials based on manganese oxide (LiMn2O4, Li(Ni1/2Mn1/2)O2, Li(Mn1/3Co1/3Ni1/3)O2), phosphates (LiFePO4), and

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silicates (Li2FeSiO4) (Tarascon et al. 1991; Padhi et al. 1997; Nytén et al. 2005). It has been shown for many of these materials that their performance is highly dependent on the synthesis method (Li et al. 2004), and indeed, the modified Pechini process has proven very useful (Liu et al. 1996). The choice of cation precursors for Li ion battery cathodes are typically nitrates and sometimes acetates (Predoana et al. 2015). Some exceptions exist, for instance, was SiO2 particles used in the synthesis of Li2FeSiO4 (giving a colloidal suspension) (Dominko et al. 2008) and tetra-n-butyl titanate was used to dope LiMn2O4 with titanium (Xiong et al. 2012). In the latter case, lactic acid was used to stabilize the titanium solution. The first report of Pechini synthesis of LiMn2O4 used EG as a solvent (Liu et al. 1996), but it has since become increasingly common to use aqueous solutions (Kunduraci and Amatucci 2006; Predoana et al. 2015). The most commonly used complexing agents are, as in Pechini’s original patent, CA and EG (Zhao et al. 2013). There are a few reports where the amount of EG has been optimized (Han and Kim 2000). Kunduraci and Amattucci prepared LiMn1.5Ni0.5O4 with a 1:1 ratio between cations and CA and an accompanying ratio of EG varying between 0, 2, 4, and 7 (Kunduraci and Amatucci 2008). Interestingly, it was found that the presence of EG had a large influence on the obtained microstructure after calcination. The materials prepared with EG had smaller particles and also a more mesoporous structure, as can be seen in Fig. 8. Both of these factors had a large impact on the electrochemical performance of the cathode material, and the optimal ratio between CA and EG was found to be 1:4. Other authors have applied only CA (Predoana et al. 2015) or a combination of CA and EDTA (Liu et al. 2014). Duncan et al. used EG without a carboxylic acid for the preparation of LiMn1.5Ni0.5O4 (Duncan et al. 2010). Here, an ester reaction occurred between EG and the acetate precursors; however, a polymer network will not form as the acetate has only one functional group. The calcination temperature has been shown to be an important parameter to optimize the electrochemical properties of the cathode materials. First of all, the

Fig. 8 Scanning electron microscopy images of LiMn1.5Ni0.5O4 prepared by a modified Pechini method with a 1:4 ratio between CA and EG (a) and without EG (b). (From Kunduraci and Amatucci 2008)

36

Modified Pechini Synthesis of Oxide Powders and Thin Films

COOH -LVP

COOH -LVP

O

OH

1107

+ CA -LVP + EG

80∞C stirring

O

COOH

Carbon shell Graphene sheet

750∞C OH - CA -LVP

N2/H2

LVP LVP

LVP

COOH-EG

Fig. 9 A schematic depiction of how a composite cathode of graphene and Li3V2(PO4)3 (LVP) was prepared from a modified Pechini process. (From Zhang et al. 2013)

volatility of Li at high temperatures is well known (Rossen et al. 1993), and sometimes an excess of Li is used in the solution to compensate for this (Predoana et al. 2015). Furthermore, a low calcination temperature will give smaller particle sizes, giving a high discharge capacity; however, the crystallinity of the particles can often be impaired, leading to fast capacity fading (Zhao et al. 2013). Typically a compromise must be found between these two aspects. A similar contradiction was encountered in the preparation of the layered material Li(Mn1/3Co1/3Ni1/3)O2 (Xia et al. 2009). Again, a low calcination temperature favorably produced smaller particles. However, in this case it also resulted in cation disorder, where Li goes into the transition metal layers, which in turn significantly reduced the lithium mobility. Finally, most of these cathode materials suffer from inherently poor electron conductivity. A carbon coating on the particle surfaces is typically necessary to mitigate this issue (Dominko et al. 2005). Sometimes this is obtained by adding a polymer, like polyethylene glycol, to the solution, followed by a heat treatment in reducing atmosphere (Mei et al. 2012). However, interestingly, when the modified Pechini method is used, this layer can be formed without the use of additives, simply from carbonization of the complexing agents CA and EG (Dominko et al. 2008). This carbon layer provides the necessary conductivity, but also serves to prevent particle growth and agglomeration (Moskon et al. 2007). Zhang et al. used a modified Pechini route to prepare a composite cathode material of Li3V2(PO4)3 and graphene (Zhang et al. 2013). Here, the complexing agents, CA and EG, were not only used to chelate the precursors, but also to attach the phosphate to functional groups on graphene oxide. In the following reducing heat treatment, CA and EG decomposed to coat the surface of the phosphate particles with a carbon layer and the graphene oxide was reduced to graphene. A schematic depiction of this process is given in Fig. 9.

Thin Films Superconducting Thin Films Superconducting thin films have been an important part of the science of superconductivity for more than six decades (Lin et al. 2015). Potential applications for hightemperature superconducting films include high-frequency electronics, microwave

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communications, and magnetic field detectors (Norton 2003). Different substrates are used for different application, and lattice match to give epitaxy and chemical compatibility are among the important considerations. Special attention has also been directed towards the preparation of superconducting wires, typically consisting of a metallic tape with a superconducting coating, which can enable large-scale applications for electric-power and magnetic applications (Larbalestier et al. 2001; Rupich et al. 2004; Kang et al. 2006). Here, a buffer layer between the metal substrate and the superconducting layer is needed to protect the metal substrate from oxidation and to give epitaxial growth of the superconducting film (Obradors et al. 2004). Chemical solution deposition has emerged as a highly competitive method for the preparation of YBCO coatings, both in the shape of wires and films, where high-quality films can be prepared cost-effectively with high speed without the need for vacuum (Obradors et al. 2006). A highly successful route for the preparation of superconducting films is the trifluoroacetate (TFA) route (Iguchi et al. 2002). The advantage of this synthesis method is that it avoids the formation of BaCO3, which can precipitate at the grain boundaries and limit the performance. However, there are some drawbacks from using fluorine, most notably the evolution of highly corrosive and dangerous hydrofluoric gas, making it less suitable for industrial upscaling (Cui et al. 2009). As a consequence, several modified Pechini routes have been developed as alternatives (Bubendorfer et al. 2003), where the formation of barium carbonate can also be avoided (Cui et al. 2009). The typical cation precursors are nitrates or acetates, but there is a larger variety in the choice of complexing agents, with CA and EG (Cui et al. 2009), EDTA (Brylewski and Przybylski 1993), triethanolamine and acetic acid (Thuy et al. 2009), and trimethylacetate and propionic acid (Shi et al. 2004) all being applied. Bubendorfer et al. investigated several organic acid, such as lactic, tartaric, glycolic, malonic, and diglycolic acid, but found that malic acid was the best choice, due to its superior ability to chelate Y3+ (Bubendorfer et al. 2003). Wang et al. used the addition of polymers, such as PVB, PEG, and PVP to improve the wettability and viscosity of the solution (Wang et al. 2008a). Solutions with a very long shelf life can be prepared (Thuy et al. 2009). By spin or dip coating on substrates such as strontium titanate or lanthanum aluminate, films with a preferential orientation can be prepared, as can be seen in the diffractogram in Fig. 10.

Transparent Conducting Oxides Powder synthesis and preparation of bulk materials, as described for ITO above, are important for the preparation of high-quality sputtering targets, but the materials are used as transparent thin films in their final application. Several physical deposition techniques for the preparation of TCOs exist, but chemical solution-based techniques offer several advantages. TCO thin films of In2O3 (Legnani et al. 2007), SnO2 (Sladkevich et al. 2011), and ZnO (Lima et al. 2007) have all been prepared by the modified Pechini method, both by spin coating (Sunde et al. 2012) and dip coating (Kundu and Biswas 2008).

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Fig. 10 XRD of an YBCO thin film with preferential orientation deposited by spin coating on a strontium titanate substrate. The solution was prepared with malic acid and glycerol as complexing agents. (From Bubendorfer et al. 2003)

The most common synthesis route is to dissolve nitrate or chloride precursors in water with the addition of CA and EG (Legnani et al. 2007). For the deposition of ZnO thin films, He et al. used a mixture of water and ethanol as solvent. In some cases, it was shown that an additive, like PVA, was necessary in order to increase the viscosity of the solution and improve the wettability on the substrate (Kundu and Biswas 2008; Sunde et al. 2014). The thickness of the films was tailored by the number of depositions, but also by varying the cation concentration and viscosity of the solution and by changing the spinning or dipping speed (Sunde et al. 2014). Layers ranging from 10 to several hundreds of nanometre could be prepared. Bernardi et al. prepared films of antimony-doped tin oxide (ATO) with varying thickness by changing the viscosity by carefully adding or evaporating water before dip coating (Bernardi et al. 2002). Two films with the same total thickness, consisting of seven and one layers, respectively, are shown in Fig. 11. They found that the highest density and best optical and electrical properties were obtained in the film made by several depositions. The TCO films are typically deposited on glass slides, but more refractory substrates, like sapphire or YSZ, are necessary if the calcination temperature is higher than about 500  C. Sladkevich et al. also prepared ATO films on sheet-like clay particles, thereby demonstrating the flexibility of the modified Pechini process (Sladkevich et al. 2011). Also patterning, which is important for TCOs used in photovoltaic devices, has been demonstrated by the modified Pechini process. Sladkevich et al. prepared films of Cd2SnO4 by spin coating of an aqueous solution with CA and EG. By using a lithographic nanoimprint technique, a pattern was made

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Fig. 11 SEM micrographs obtained at 45 inclination of ATO films prepared by dip coating of a modified Pechini solution. The top film (a) is prepared by seven layers from a solution with a viscosity of 4 cP. The bottom film (b) is a single layer from a solution with 20 cP. (From Bernardi et al. 2002)

in the as-spun polymeric precursor film, which remained in the film after calcination (Schumm et al. 2011). Here, it was also demonstrated that they could obtain the metastable cubic phase of Cd2SnO4 in the thin film, which cannot be obtained in bulk materials, by optimizing the calcination procedure. Finally, the calcination temperature and atmosphere have been shown to be very important for the electrical properties of the TCOs (Choppali and Gorman 2008). After heat treatment at high temperatures and in reducing atmospheres, ITO thin films with excellent properties have been obtained. With a specific resistance down in the 10 4 Ω
cm range, the films prepared by the modified Pechini method are comparable to the best values from physical deposition techniques (Sunde et al. 2014).

Thin Film Phosphors Thin films of luminescent materials are attractive for many technological applications, especially related to display technology (Yu et al. 2005), but also for lightconversion layers for photovoltaics (Huang et al. 2013) and optical waveguides (Chae et al. 2013). Already in 1980 Robertson and van Tol demonstrated that epitaxial luminescent films of rare earth-doped garnets could withstand much higher power densities in cathode ray tubes without degradation than their powder counterparts (Robertson and Van Tol 1980). Since then, thin film phosphor materials have received significant attention (Choe et al. 2001; Garskaite et al. 2010). The uniform thickness and smoother surface of the thin films makes it possible to define smaller pixels, thereby giving a higher resolution. Thin films prepared by the modified Pechini method can also be patterned, which can be achieved by relatively simple

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Fig. 12 Optical photographs of patterned thin films of LaPO4 doped with Ce3+ and Tb3+ prepared by the modified Pechini process and the micromolding in capillaries technique. (Modified from (Lin et al. 2007). Reprinted with permission from Lin et al. Copyright 2007 American Chemical Society)

and inexpensive soft-lithography techniques (Pang et al. 2003). A demonstration of such patterning is given in Fig. 12. These patterns were prepared by the micromolding in capillaries technique. Here, a droplet of the modified Pechini solution, containing the cation precursors and complexing agents, is deposited next to a micromold, upon which capillary forces will pull the solution into the mold. After drying the mold can be removed, and the patterned oxide thin film remains after calcination (Lin et al. 2007). Phosphors made up from monodisperse, small, and spherical particles are attractive due to the possibility of a low light-scattering and a high packing density, leading to good resolution (Martinez-Rubio et al. 2001). In this regard, the flexibility of the modified Pechini method can be utilized to produce core-shell structures. Here, monodisperse and spherical silica particles are typically produced by the Stöber method (Stöber et al. 1968). These particles are added to a modified Pechini solution, containing the chosen cations in the desired stoichiometry together with complexing agents, typically CA and PEG (Wang et al. 2005). The silica particles contain OH groups, which the chelated cations can bond to, thereby producing a thin coating after calcination. By using a silica core, the total cost of the phosphor particle is significantly reduced. The optical properties of the core-shell particles can be tuned by the number of coatings, the calcination temperature, and the size of the silica particle. The crystallinity of the coating improves with increasing annealing temperature; however, if the temperature is too high, a reaction between the core and the shell can occur (Lin et al. 2007).

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Smet PF, Parmentier AB, Poelman D. Selecting conversion phosphors for white light-emitting diodes. J Electrochem Soc. 2011;158:R37–54. Stöber W, Fink A, Bohn E. Controlled growth of monodisperse silica spheres in the micron size range. J Colloid Interface Sci. 1968;26:62–9. Sunde TOL, et al. Transparent and conducting ITO thin films by spin coating of an aqueous precursor solution. J Mater Chem. 2012;22:15740–9. Sunde TOL, Einarsrud MA, Grande T. Solid state sintering of nano-crystalline indium tin oxide. J Eur Ceram Soc. 2013;33:565–74. Sunde TOL, Einarsrud MA, Grande T. Optimisation of chemical solution deposition of indium tin oxide thin films. Thin Solid Films. 2014;573:48–55. Tai LW, Lessing PA. Modified resin – intermediate processing of perovskite powders: Part I. Optimization of polymeric precursors. J Mater Res. 1992a;7:502–10. Tai LW, Lessing PA. Modified resin – intermediate processing of perovskite powders: Part II. Processing for fine, nonagglomerated Sr-doped lanthanum chromite powders. J Mater Res. 1992b;7:511–9. Tao S, Irvine JTS. Synthesis and characterization of (La0.75Sr0.25)Cr0.5Mn0.5O3-δ, a redox-stable, efficient perovskite anode for SOFCs. J Electrochem Soc. 2004;151:A252–9. Tarascon JM, Wang E, Shokoohi FK, McKinnon WR, Colson S. Spinel phase of LiMn2O4 as a cathode in secondary lithium cells. J Electrochem Soc. 1991;138:2859–64. The IUPAC Stability Constants Database, Academic Software. http://www.acadsoft.co.uk/ Thuy TT, et al. Sol–gel chemistry of an aqueous precursor solution for YBCO thin films. J Sol–Gel Sci Technol. 2009;52:124–33. Van der Biest O, et al. Ceramic superconductors synthesized by sol–gel methods. Phys C. 1991;190:119–21. Vojnovich T, Bratton RJ. Impurity effects on sintering and electrical resistivity of indium oxide. Am Ceram Soc Bull. 1975;54:216–7. Wang J, et al. Epitaxial BiFeO3 multiferroic thin film heterostructures. Science. 2003;299:1719–22. Wang H, Lin CK, Liu XM, Lin J, Yu M. Monodisperse spherical core-shell-structured phosphors obtained by functionalization of silica spheres with Y2O3:Eu3+ layers for field emission displays. Appl Phys Lett. 2005;87:1–3. Wang WT, et al. Chemical solution deposition of YBCO thin film by different polymer additives. Phys C. 2008a;468:1563–6. Wang Z, et al. NaEu0.96Sm0.04(MoO4)2 as a promising red-emitting phosphor for LED solid-state lighting prepared by the Pechini process. J Lumin. 2008b;128:147–54. Wang H, Xu X, Li X, Zhang J, Li C. Synthesis and sintering of indium tin oxide nanoparticles by citrate-nitrate combustion method. Rare Met. 2010;29:355–60. Wu MK, et al. Superconductivity at 93 K in a new mixed-phase Y-Ba-Cu-O compound system at ambient pressure. Phys Rev Lett. 1987;58:908–10. Xia H, Wang H, Xiao W, Lu L, Lai MO. Properties of LiNi1/3Co1/3Mn1/3O2 cathode material synthesized by a modified Pechini method for high-power lithium-ion batteries. J Alloys Compd. 2009;480:696–701. Xiong L, Xu Y, Tao T, Goodenough JB. Synthesis and electrochemical characterization of multi-cations doped spinel LiMn2O4 used for lithium ion batteries. J Power Sources. 2012;199:214–9. Yang YM, et al. Characterization of YBa2Cu3O7-x bulk samples prepared by citrate synthesis and solid-state reaction. J Appl Phys. 1989;66:312–5. Ye S, Li Y, Yu D, Yang Z, Zhang Q. Structural effects on Stokes and anti-Stokes luminescence of double-perovskite (Ba, Sr)2CaMoO6: Yb3+, Eu3+. J Appl Phys. 2011;110:013517. Yu M, Lin J, Wang SB. Effects of x and R3+ on the luminescent properties of Eu3+ in nanocrystalline YVxP1-xO4:Eu3+ and RVO4:Eu3+ thin-film phosphors. Appl Phys Mater Sci Process. 2005;80:353–60. Zalga A, Moravec Z, Pinkas J, Kareiva A. On the sol–gel preparation of different tungstates and molybdates. J Therm Anal Calorim. 2011;105:3–11.

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Zhang H, Fu X, Niu S, Xin Q. Synthesis and photoluminescence properties of Eu-doped AZrO3 (A = Ca, Sr, Ba) perovskite. J Alloys Compd. 2008;459:103–6. Zhang L, et al. Li3V2(PO4)3@C/graphene composite with improved cycling performance as cathode material for lithium-ion batteries. Electrochim Acta. 2013;91:108–13. Zhao Y, Sun G, Wu R. Synthesis of nanosized Fe-Mn based Li-rich cathode materials for lithiumion battery via a simple method. Electrochim Acta. 2013;96:291–7. Zhu WZ, Deevi SC. Development of interconnect materials for solid oxide fuel cells. Mater Sci Eng A. 2003;348:227–43.

Part II Characterization

Characterization of Sol-Gel Materials by Infrared Spectroscopy

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Rui M. Almeida and Ana C. Marques

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Material Preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Infrared Spectroscopic Techniques . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Silica Sol-Gel Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conversion of Alkoxide Precursors to Silica Gel and Glass . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Bulk and Thin-Film Gels and Gel-Derived Glasses . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . TO and LO Vibrational Mode Splitting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . OH Species . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Relationship Between Infrared Absorption and Porosity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Inorganic Versus Hybrid SiO2 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Multicomponent Silicates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Structural Evolution with the Heat Treatment, Temperature and Time . . . . . . . . . . . . . . . . . . . . Evolution of the TO Peak Frequency . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Comparison Between Different Systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Influence of the Thickness and Number of Layers on the Microstructure of Multilayer Silicate Films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nonsilicates and Non-oxides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1122 1123 1123 1126 1126 1127 1128 1130 1131 1132 1134 1137 1137 1138 1140 1142 1144 1148 1148

R. M. Almeida (*) Departamento de Engenharia Química / Centro de Química Estrutural (CQE), Instituto Superior Técnico / Universidade de Lisboa, Lisboa, Portugal e-mail: [email protected] A. C. Marques Departamento de Engenharia Química / Centre for Natural Resources and the Environment (CERENA), Instituto Superior Técnico, Universidade de Lisboa, Lisboa, Portugal e-mail: [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_33

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Abstract

Infrared (IR) spectroscopy has been widely used in the characterization of sol-gel (SG) materials, leading to an extensive knowledge of the relationship between IR spectra, atomic level structure and properties of bulk materials, and mainly thin films. Moreover, the development of in situ time-resolved IR spectroscopy has enabled the monitoring of structural changes during physical phenomena occurring in the advanced processing of SG-derived materials. This chapter reviews the various IR spectroscopy techniques which have been used in the study of SG materials, including silica-based inorganic and hybrid gels plus other oxide and non-oxide compositions, with particular emphasis on the relationships between the IR spectra and the presence of OH species, porosity, longitudinal and transverse vibrational mode splittings, and the general problems involved in the densification of SG-derived coatings and bulk materials.

Introduction Infrared spectroscopy is generally a nondestructive technique, which provides extensive information about the optical, vibrational, chemical, and structural properties of many different classes of materials and, in particular, of SG materials. The fact that many SG materials are prepared as coatings, in thin-film form, is an added reason why IR spectroscopy is ideally suited for the characterization of these materials. It is a highly adequate technique for the analysis of silicates, as well as other oxides and non-oxides obtained by the SG method. IR spectroscopy can be performed on bulk SG samples, as well as on thin films, usually prepared by spin or dip coating, which are the most common coating deposition techniques. In the latter case, an IR transparent substrate like silicon is often used. IR spectroscopy, nowadays performed almost exclusively on Fourier transform IR (FTIR) spectrometers, is a highly versatile technique, and depending on the different types of samples, several methods can be used, including absorption spectroscopy, specular reflectance, diffuse reflectance IR Fourier transform spectroscopy (DRIFTS), attenuated total reflectivity (ATR), IR reflection/absorption spectroscopy (IRRAS), microspectroscopy, and photoacoustic spectroscopy (PAS) (Almeida 1988; Almeida et al. 1990a; Stuart et al. 1996; Chalmers and Dent 1997). The term “spectroscopy” strictly refers to double-beam instruments, whereas “spectrometry” refers to spectra taken with single-beam instruments, which form the majority of the present-day FTIR instruments; the two terms will be used interchangeably throughout this chapter. IR spectroscopy, the most common form of vibrational spectroscopy, together with Raman spectroscopy, has been one of the most utilized tools to investigate the sol-to-gel evolution and gel aging in SiO2-based materials, the microstructural evolution of bulk and thin-film gel materials with temperature, including their thermal densification (with the elimination of porosity and OH species), the occurrence of phase separation, and the influence of processing parameters on the microstructure. Silica-based materials, in particular, have been the most extensively studied SG materials.

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

1123

Among the theoretical aspects involved in the use of IR spectroscopy, the vibrational activity is the most relevant one. In fact, some vibrational modes have stronger activity in the IR than others and this determines the corresponding peak absorption intensity. The IR activity of a certain fundamental vibrational mode depends on the general IR selection rule which states that the permanent dipole moment of a given molecular group (μ) must have a nonzero derivative, with respect to the particular normal coordinate, for the equilibrium configuration of the molecular group in question. Thus, a vibration must cause a change in the corresponding dipole moment, in order to give rise to absorption of IR radiation. The larger this change, the more intense the absorption will be (Nakamoto 1978; Almeida 1987). In this chapter, particular attention will be devoted to the study of the structure of SG-derived silica and modified silicates, at the molecular level, by IR absorption and reflection spectroscopies, as well as to microstructural aspects such as the elimination of residual porosity during thermal densification, which usually occurs together with a simultaneous elimination of residual OH species. Relevant results for other nonsilicate and non-oxide materials will also be reviewed.

Material Preparation SG chemistry can be catalyzed by acids or bases, the former being the most common case. In acidic media, the starting point of the preparation is the hydrolysis of a mixture of alkoxide precursors, with water, in solvent such as ethanol (EtOH), or other alcohols. This is followed by condensation (polymerization) and sol formation. This colloidal solution is then aged and the evaporation of the solvent and water may occur slowly, leading to a bulk gel, or rapidly, by deposition methods such as spin or dip coating. The porous bulk or thin-film gel thus obtained may be further dried and sintered (densified), with simultaneous elimination of porosity and residual OH groups (Seco et al. 2000). SG can also be combined with colloidal chemistry to introduce a controlled structuration of the final materials at the mesoscale. For more details on the preparation of SG materials, the reader is referred to Vol. 1 of this series of books or to a general treatise like that by Brinker and Scherrer (1990).

Infrared Spectroscopic Techniques IR absorption or reflection spectroscopies are standard techniques of optical vibrational characterization, where IR radiation is used. The IR spectrum can be generally divided into three frequency (or wave number/wavelength) regions: the far IR (10–400 cm
1), the middle IR (400–4000 cm–1), and the near IR (2500–700 nm), with the middle IR region being employed in most IR spectroscopic investigations. IR absorption spectroscopy is based on the absorption of IR radiation at frequencies corresponding to the vibrational modes of molecular groups or crystals, as it passes through a material sample, in solid or liquid form. The radiation beam usually strikes the sample at normal incidence and passes through the sample into the

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spectrometer detector, unless absorbed by the sample, or lost by other processes, such as scattering or back reflection from the surfaces (Stuart et al. 1996; Chalmers and Dent 1997). For thin films, the first-order IR absorption spectra may be obtained when their thickness is approximately within one order of magnitude of the wavelength of the IR radiation used. The absorbance, A, is proportional to the absorption coefficient, α, as given by Beer’s law: A ¼ αx ¼ log10

  1 T

(1)

where x is the film thickness and T is the transmittance, T ¼ II0 ¼ expð
α0 xÞ, with α0 = 2.303α. Bulk samples may be ground, mixed with an IR transparent powder such as KBr (for the middle IR)
or polyethylene (for far IR), and pelletized. Here, α ¼ 4πλ k, where k is the dimensionless extinction coefficient, which is the imaginary part of the complex refractive index, n* = n – ik (n being the real part of the refractive index). IR absorption spectroscopy can also be performed with IR radiation at oblique incidence. The advantage of using oblique incidence is to reveal structural differences that would otherwise remain hidden in the conventional, normal incidence spectra, some examples of which are given in a latter section. IR reflection spectroscopy is normally used for samples which are difficult to analyze in transmission, such as bulk samples or thin layers on nontransparent substrates. In this case, the IR radiation is directed at a sample surface, usually at an angle larger than 0 off-normal, with the attenuated radiation, reflected back from that surface, being detected. Reflection techniques can be based on specular reflection (internal or external reflection), where the reflectivity R, at normal incidence, is given by R¼

ð n
1Þ 2 þ k 2 ð n þ 1Þ 2 þ k 2

;

(2)

and thus depends on the absorption (k) and refraction (n). Attenuated total reflectance (ATR) spectroscopy is based on the phenomenon of total internal reflection. The evanescent field of a beam of radiation penetrates a fraction of a wavelength beyond the reflecting surface of an optically dense dielectric waveguide, and when a material which selectively absorbs radiation is in close contact with the reflecting surface, the beam loses energy at the wavelengths where the material absorbs. The depth of penetration is a function of the wavelength, the refractive index of the material, and the angle of incident radiation (Stuart et al. 1996). IR reflection–absorption spectroscopy (IRRAS) utilizes the phenomenon of external reflection; reflection–absorption spectra are usually collected with polarized IR light, at a certain angle off-normal to the surface. It is possible, by this method, to study the spectroscopic behavior for the two light polarizations: s (electric field perpendicular to the plane formed by the incident and reflected rays) and p (electric field of the radiation parallel to the same plane) (Almeida and Pantano 1990b). Figure 1 shows

Characterization of Sol-Gel Materials by Infrared Spectroscopy

Fig. 1 IR reflection–absorption (IRRAS) spectra, at 55 off-normal incidence, of a silica gel film dried at room temperature, taken with (s), perpendicularly polarized radiation; (p), parallelpolarized radiation (Reproduced with permission from Almeida and Pantano (1990b))

1125 1075

1075

% TRANSMITTANCE

37

P

s 1232 4000

3000

2000

1000

FREQUENCY (CM−1)

an example of spectra collected by this method for a silica gel film, where the s-spectrum shows the transverse optical (TO) component of the dominant Si–O–Si asymmetric stretching vibration, at 1075 cm
1, whereas the p-spectrum reveals the longitudinal optical (LO) component of the same mode, at 1232 cm–1, in addition to the TO peak. The physical meaning of the TO and LO spectra will be examined in detail in “TO and LO Vibrational Mode Splitting.” Another widely used technique, especially in the case of powders, is DRIFTS, where the IR radiation impinges on the sample over a range of angles simultaneously, for instance, between 20 and 70 off-normal, and is collected by parabolic mirrors over a large solid angle. An advantage of using DRIFTS is the ease of sample preparation, where powdered or bulk specimens can be used, the latter without the need of being polished (Almeida et al. 1990a). A more recently developed technique is time-resolved IR spectroscopy, which is based on standard IR absorption spectroscopy, but consists of rapid-scan timeresolved measurements, performed by averaging a number of interferograms per spectrum within a predetermined acquisition time and using a selected time interval between the consecutive spectra. This is a very good experimental tool to follow time-dependent phenomena with good temporal resolution. In particular, it has been employed for gathering information about the chemistry of SG/colloidal chemistry processes such as evaporation-induced self-assembly (EISA), namely, the polycondensation reactions that occur during the transition from sol to gel of the evaporating liquid layer. For such purpose, it has been combined with grazing incidence smallangle X-ray scattering (GISAXS, using synchrotron radiation), as reported in an exploratory work by Marcelli et al. (2012).

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Silica Sol-Gel Materials Conversion of Alkoxide Precursors to Silica Gel and Glass The different stages of SG processing of silica gels have been the subject of several IR studies (Matos et al. 1992; Niznansky and Rehspringer 1995; Gnado et al. 1996). Those stages include the formation of oligomeric species during the hydrolysis/ polycondensation of tetraethoxysilane (TEOS, the most widely used silica precursor) and the TEOS to silica gel and glass conversion. Figure 2 shows sequences of IR spectra, recorded as the hydrolysis and condensation proceeded, in the system TEOS–EtOH–H2O–HCl (molar ratio of TEOS: EtOH:H2O  1:6:6). The main changes that may be observed, as the reactions proceeded, consist in the disappearance of the TEOS peaks at 473, 1102, 1168, 1299, and 1400 cm–1 (Matos et al. 1992; Vis et al. 1992), some of them already hardly visible in Fig. 2b, after 130 min. In addition, one may observe an increase in the band at 881 cm
1, due to CH3 or CH2 deformation in EtOH, during the first 7 h of the sol-to-gel conversion,

960 881

1085

584 793

b

1048

473 434

a 1168

c

954

Transmittance

Fig. 2 IR spectra taken during the TEOS to silica gel conversion. The reaction time after mixing of the component liquids was (a) 30 min, (b) 2 h 10 min, (c) 5 h 30 min, (d ) 7 h 10 min, (e) 8 h, and ( f ) 28 h. The bands marked with (*) are due to atmospheric CO2 (Reprinted from Journal of Non-Crystalline Solids, 147 & 148, M.C. Matos, L.M. Ilharco, R.M. Almeida, “The evolution of TEOS to silica gel and glass by vibrational spectroscopy”, p. 234, Copyright (1992), with permission from Elsevier)

d

1087

e

f 799 960 −460

1163

1078

1500

950 Wavenumber / cm−1

∗

400

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

1127

followed by a substantial reduction of its intensity, afterward. This is an evidence for the formation of ethanol in the initial steps of hydrolysis. Water-forming condensation, which starts while hydrolysis is still in progress, will cause the elimination of ethanol, plus the oligomer cross-linking. Broad and strong absorption bands are observed between 1000 and 1250 cm
1 and near 950 cm
1. Their positions, intensities, and shapes also change during gel aging. The relatively sharp band at 960 cm
1 (rocking of CH3 in TEOS) is replaced by a somewhat broader band near the same position, characteristic of a wet gel (indicating the presence of Si–OH or Si–O– species). Figure 2f is dominated by the three main features of the IR spectrum of a silica xerogel, due to Si–O–Si bonds: a new band at 1078 cm
1, which presents a clearly visible high-frequency shoulder at 1163 cm
1 and a peak at 460 cm
1, stronger than the peak at 434 cm
1 in the initial stages of the process. In summary, the sol-to-gel evolution is based on the ethanol and TEOS consumption and on the formation of Si–O–Si bonds, which is demonstrated by the disappearance of the ethanol and TEOS-related bands and by the simultaneous formation and intensification of the siloxane (Si–O–Si)-related bands. Finally, the gel-to-glass conversion may be followed through the disappearance of the IR peaks due to unhydrolyzed alkoxy groups and to Si–OH stretching, together with the intensification of the peaks at 800 and 1078 cm
1. This is suggested by Matos et al. (1992) and occurs for heat treatments above 700  C, indicating that the films are close to full densification after such heat treatments. Despite the fact that acid catalysts are the most employed in SG processing, fresh and heat-treated gels, obtained via basic catalysis, have also been analyzed by IR absorption spectroscopy. In particular, Popescu et al. (1995) investigated and characterized quantitatively the intermediate-range order in silica gels, in particular its smearing in fresh gels with respect to v-SiO2 and its recovery upon heat treatment. IR spectroscopy has been, therefore, very useful in the study of the influence of the type of catalyst used in the processing of SG materials.

Bulk and Thin-Film Gels and Gel-Derived Glasses SG processing allows the preparation of bulk glasses without going through the melting process. Namely, the preparation of monolithic gel pieces may be achieved by the slow drying of xerogels, directly molded from SG solutions (Thomas 1988), although great caution and slow heat treatments are needed in order to prevent cracking from occurring. Aggregation, gelation, and drying (with the associated shrinkage) usually occur in days or weeks, for bulk materials. On the other hand, in thin-film deposition, the time spent on these processes consists only of some seconds or minutes. This results in a lower cross-linking degree, leading to a somewhat more compact structure in films, compared with bulk gels (Fardad et al. 1995), which is confirmed by the fact that the refractive index of silica glass films at 633 nm is slightly higher than that of bulk vitreous SiO2 (Almeida and Pantano 1990a). Due to this difference, IR studies of the bulk may be not entirely appropriate for interpreting structural information of films. Also, the film deposition method (dip and spin

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coating being the most common ones) has some effect on the densification process. The spin-coating process, especially for large spinning frequencies, leads to denser films, possibly because of the higher evaporation rate, relative to the condensation rate (Almeida 1994). All the above phenomena may be assessed by IR spectroscopy. For silicates, two of the main features associated with higher densification levels are the increase in frequency and intensity of the Si–O–Si stretching peak at 1070 cm
1 and the simultaneous reduction in intensity of the OH-related peaks.

Structure Figure 3 shows the IR absorption spectrum (in transmission mode) of a free-standing silica gel foil, about 1 μm thick, dried at room temperature. It exhibits a strong O–H stretching peak at 3400 cm
1, with some free OH species (high-frequency shoulder), a few free surface silanol (Si–OH) groups at 3740 cm
1, and a majority of hydrogen-bonded OH groups (Almeida and Pantano 1990a). It also shows small features near 2900–3000 cm
1, due to organic residues (ethanol and TEOS). In addition, a very weak peak at 1650 cm
1 can be observed, which could in part be due to residual ethanol, indicating that very little molecular water (the other possible vibrating species at 1650 cm
1), if any, remained in the gel foil at this stage. Figure 4 shows the IR absorption spectra, in the fundamental Si–O region, of the same free-standing gel foil of Fig. 3 and a gel film spun on a single crystal Si substrate, both dried at room temperature. Comparing the spectra, both exhibit bands centered near 933–950 and 578 cm
1. The higher-frequency peak may be associated with stretching vibrations of Si–OH or Si–O– groups (Almeida et al. 1990b) and the

100.0

TRANSMITTANCE (%)

Fig. 3 IR absorption spectrum of a free-standing silica gel foil, about 1 μm thick (Reproduced with permission from Almeida and Pantano (1990a))

50.0

0.0 4000

3000

2000

FREQUENCY

1000

[cm−1

]

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

ABSORBANCE UNITS

1.5

1129

1075

0.75

1070

950

933

0.0 1400

1200

1000

800

600

400

−1

FREQUENCY [cm ]

Fig. 4 IR absorption spectra of (—) an undensified silica gel film and (----) a silica gel foil (Reproduced with permission from Almeida and Pantano (1990a))

lower-frequency peak may be attributed to rocking motions of the same species (Almeida and Pantano 1990a). The other three main features are assigned to different vibrational modes of the Si–O–Si bonds, which can be taken as the silica fingerprints. They are the high-frequency dominant band near 1070 cm
1, the weak intermediate-frequency band at 800 cm
1, and the lower-frequency band near 450 cm
1. The latter one is due to rocking motions of the oxygen atoms perpendicular to the Si–O–Si plane, accompanied by some Si cation motion. The one at 800 cm
1 may be assigned to symmetric stretching of the oxygen atoms along the bisector of the Si–O–Si bridging angle, again with some simultaneous silicon cation motion (Galeener 1979; Almeida and Pantano 1990a). The peak at 1070 cm
1, which will be discussed in greater detail in the next section, is due to an asymmetric stretch of the oxygen atoms. At this point, one can attempt to interpret this peak in terms of structure, based on the central force network model of Sen and Thorpe (1977) and Galeener (1979). Almeida and Pantano (1990a) used this model to interpret the FTIR data of SG silica films, assuming a short-range order with a unique Si–O bond distance, characterized by a single force constant value, k, and a distribution of Si–O–Si intertetrahedral angles, θ. The frequency of the Si–O–Si asymmetric stretch is given by ω2 ¼

k 4 k ð1
cos θÞ þ m0 3 mSi

(3)

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R. M. Almeida and A. C. Marques

Differentiating this equation with respect to θ, one obtains dω k sin θ ¼ dθ 2ωm0

(4)

Rearranging this expression, one has the following equation for the spread, Δθ, in the intertetrahedral angle, as a function of the full width at half maximum (FWHM), Δv, of the peak: Δθ ¼

vΔv ð530:5k sin θÞ

(5)

An increase in the FWHM of the peak near 1070 cm
1 corresponds to a broadening of the bond angle distribution. The values of 24 , 27 , and 30 were found for the width of the bond angle distributions, in room temperature dried, 400  C treated and glassy SiO2 films, respectively (Almeida and Pantano 1990a).

TO and LO Vibrational Mode Splitting In this section, particular attention will be devoted to the peak of SiO2 at 1070 cm
1 and its high-frequency shoulder. This peak has been used to obtain information concerning strains on the Si–O–Si bonds (related to the normal values of the intertetrahedral angle and bond length) (Almeida and Pantano 1990a), porosity, and thickness of silica films (Almeida et al. 1994). The high-frequency, dominant band near 1070 cm
1, in Figs. 3 and 4, is due to the transverse optical, or TO, component of the asymmetric stretching of Si–O–Si bonding sequences (Almeida and Pantano 1990a), accompanied by Si cation motion. The shoulder (Sh) at 1200 cm
1 has been shown to be related to the longitudinal optical (LO) component of the same vibration (Almeida and Pantano 1990a). For films which are thin compared to the IR wavelengths, radiation incident normal to the surface should only excite TO modes, whereas radiation oblique to the surface should excite both TO and LO modes (Almeida 1992). However, as experimentally evidenced (see Fig. 4), the Sh band, associated with the LO mode, appears also at normal incidence, because there probably are no pure LO and TO modes in the glass, due to the lack of long-range order (Almeida et al. 1990a). The intensification of the Sh band in incompletely densified films (Fig. 4) may be associated with the presence of residual film porosity, which scatters a normal incidence IR beam in all directions (Almeida and Pantano 1990a) and leads to the activation of the LO mode. Therefore, the Sh/TO intensity ratio can be used as a semiquantitative measure of the residual porosity of the gel film, a function of its degree of densification. The major difference noticed between the undensified silica gel film and the silica gel foil spectra (Fig. 4) was the much higher value of the Sh/TO intensity ratio for the thicker gel foil. This may suggest that the latter, being more porous, scatters

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

1131

the 10 μm wavelength light to a larger extent, sending it in all directions, even for light originally at normal incidence. The fraction of light which effectively becomes off-normal incident in this way will further activate the LO mode, and this effect increases with porosity (as observed), as well as with the film thickness, up to the penetration depth of the IR radiation (1/nα). Despite the extensive number of works based on the subject, there is still a certain disagreement concerning the IR band assignments for silica. The high-frequency shoulder (the Sh band at 1200 cm
1) has also been proposed by some authors, based on peak fittings, to evolve at oblique incidence into two separate bands of LO character: a 1250 cm
1 peak, due to the gel skeleton, and another band near 1160 cm
1, assigned to strained Si–O–Si bridges at the surface of the gel film pores (Gallardo et al. 2002; Innocenzi 2003). On the other hand, some authors assign the 1000–1260 cm
1 region bands, especially the Sh band, to cyclic species, in particular to strained structural rings (fourfold siloxane rings) (Fidalgo and Ilharco 2001). These bands have been observed in SG silica films prepared with low H2O/ alkoxide ratios and often also in hybrid silica materials, but they disappear for heat treatment temperatures above 300  C (Innocenzi 2003). The identification of cyclic species and disorder-induced modes in the 1000–1260 cm
1 region still needs, however, a good deal of basic research in order to be fully clarified.

OH Species IR spectroscopy is not only able to detect easily the presence of OH-related species, but it can also provide much information on molecular interactions involving hydrogen bonding and it has the advantage of being direct and very sensitive. The hydrogen bond most clearly recognized and easily studied is the intermolecular bond (Chalmers and Dent 1997), whereas the intramolecular bond is less easy to detect. The intramolecular hydrogen bond related to OH peak is broader and it has a lower cross section than both the free OH and the intermolecular hydrogen bond peaks. Structural hydroxyl groups in glass have fundamental O–H stretching vibrations between 2700 and 4000 cm
1, depending on the position of those groups in the glass network (Almeida and Pantano 1990a). Monoliths and thin films, both dried at room temperature or at higher temperatures, may contain (Fidalgo and Ilharco 2001; Innocenzi 2003) a residual content of: – Molecular adsorbed water, with peaks at 3300–3500 cm
1, assigned to O–H stretching in H-bonded water and also at 1650 cm
1, due to bending vibrations of the water molecules – Ethanol, with peaks at 3300–3650 cm
1, due to H-bonded OH vibrations of the alcohol molecules and also at 880 cm
1, due to methyl and methylene vibrations – Silanol groups (Si–OH), containing isolated (or free) OH species, responsible for a sharp peak at 3740 cm
1 and terminal OH species, with a stretching mode in the region of 3600–3800 cm
1), as well as hydrogen-bonded OH groups (stretching at 3650–3200 cm
1), as shown in Fig. 5

1132

R. M. Almeida and A. C. Marques

a

Terminal

Isolated H

H

H

O Si

H

O O

O

Si

Si

Terminal

H2O

b H

H

H

H

H

O

O

O

O

O

Si

Si

Si

Si

Si

H O Si

O Si

Si

Hydroxyl chain Fig. 5 (a) Schematic example of isolated and terminal hydroxyl groups. (b) H-bonded hydroxyl chain: example of silanol condensation with elimination of water and formation of isolated hydroxyls on the silica film surfaces (Reprinted from Journal of Non-Crystalline Solids, 316, P. Innocenzi, “Infrared spectroscopy of sol-gel derived silica-based films: a spectramicrostructure overview”, p. 318, Copyright (2003), with permission from Elsevier)

There is a strong need for the use of ultralow H2O-containing glass, namely, for highly transparent glasses, for optical and electronic applications. IR spectroscopy allows the assessment of the OH content evolution in the glass network with the thermal treatment of the precursor gels, for heat treatment atmospheres of different humidity, as well as with different processing parameters for the sol preparation. In fact, the bands related to OH stretching modes decrease in intensity when the heat treatment temperature increases or when special chemical treatments with reactive gases, such as CCl4 or Cl2 (Mukherjee 1988), are performed on the samples. H-bonded silanols, which are believed to be located at the pore surfaces, can form chains of different lengths and shapes, giving rise to a corresponding broad band at 3500–3600 cm
1. Upon heat treatment, there is elimination of water, a decrease in the number of those chains, with the formation of siloxane bridges, and an increase of isolated species. Therefore, the decrease in intensity of the broad band due to chain silanols occurs generally with an increase of the 3740 cm
1 peak (Innocenzi 2003).

Relationship Between Infrared Absorption and Porosity This is a relationship which can be applied to the estimate of the residual porosity of SG films. Their volume fraction of porosity, Vp, can be estimated from the film thickness, obtained, for example, by ellipsometry or mechanical profilometry and from the so-called IR thickness (Almeida et al. 1994), determined from their IR

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

1133

spectra. The IR thickness of the porous film is the thickness which it will exhibit after full densification; this thickness, which is independent of the initial film porosity, is always less than the one measured for the porous film. As densification occurs at increasing temperatures, the IR thickness remains constant, but the experimentally determined, profilometer (or ellipsometer) thickness decreases (Almeida et al. 1994). This is due to the fact that the volume percent porosity decreases with the temperature and time of heat treatment, but the amount of solid material remains basically the same. Defining “IR thickness” of a material as the thickness xIR that it would have after full densification, therefore independent of the actual processing method and residual porosity, Beer’s law may be written in the form A ¼ α  x ¼ αstd  xIR

(6)

where A is the measured absorbance of the film, x is its measured thickness (in cm), α is the absorption coefficient of the porous film (in cm
1), xIR is the “IR thickness” (or the effective thickness of the fully densified, absorbing solid), and αstd is the absorption coefficient of a densified, standard film. Thus, the IR thickness is defined as xIR ¼ A=αstd

(7)

If the film has a certain volume fraction of porosity Vp, xIR is related to the measured thickness x by xIR ¼ k  x

(8)

where k = 1 – Vp = α/αstd. Consequently, Vp will be given by 

α Vp ¼ 1
αstd

 ¼1


x  IR x

(9)

Therefore, simply by knowing the values of α (from IR spectra and thickness measurements) and αstd, one can estimate the porosity (Vp) of a film. Figure 6 shows plots of absorbance versus thickness, which enable the determination of the coefficient α = k  αstd, from linear regression fittings to Beer’s law. The samples studied consisted of 100% SiO2 and 80% SiO2–20% TiO2 (molar percentages are used throughout this chapter), subjected to thermal treatments at 550  C and 450  C, respectively. This temperature has been chosen, because it is sufficiently high for multilayer deposition without cracking and there is still a significant residual Vp. By using αstd values, for fully densified films, of 1.4  104 cm
1, for silica and 1.0  104 cm
1, and for silica–titania, the values obtained for k were 0.84 and 0.71, respectively. This corresponds to a volume percent porosity of 16%, for pure silica films densified at 550  C and to a volume percent porosity of 29%, for silica–titania films densified at 450  C.

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R. M. Almeida and A. C. Marques

3.0

2.5

Absorbance

2.0

Slope 0.0011−

1.5

1.0 Slope 0.00071

0.5 100 S 80S20T

0.0 0

500

1000

1500

2000

2500

Thickness / nm

Fig. 6 Absorbance of the Si–O–Si band at 1070 cm–1 versus thickness, for silica (100% SiO2) and 80% SiO2–20% TiO2 multilayer films (Reproduced with permission from Almeida et al. (1994))

Inorganic Versus Hybrid SiO2 IR spectroscopy provides also information on the structural and compositional properties of hybrid organic/inorganic SG networks. Hybrid SG materials, for example, may be prepared by incorporating a certain amount of an organically modified precursor, such as methyl triethoxysilane, MTES, containing non-hydrolyzable Si–CH3 groups, into the initial solution. Several works have been performed with the aim of studying and comparing the structural evolution of inorganic and hybrid silica SG coatings during sintering (Innocenzi et al. 1994; Seco et al. 2000; Fidalgo and Ilharco 2001; Almeida et al. 2002; Gallardo et al. 2002). The evolution of the normal incidence IR spectra of inorganic silica (IS) and hybrid silica (HS) films with the heat treatment temperature is shown in Fig. 7. In addition to the main peak characteristic of the vibrational modes of Si–O–Si bonds, common to both spectra, the spectra related to the hybrid films (molar ratio of TEOS/ MTES = 2:3) present some new features, when compared to the inorganic silica films. A small peak at 1275 cm
1, corresponding to the C–H vibrations in Si–CH3

Characterization of Sol-Gel Materials by Infrared Spectroscopy

Fig. 7 Evolution of FTIR spectra for coatings prepared from TEOS (IS) and MTES/ TEOS (HS) at normal incidence, for different heat treatment temperatures (Reprinted from Journal of Non-Crystalline Solids, 298, J. Gallardo, A. Durán, D. Di Martino, R.M. Almeida, “Structure of inorganic and hybrid SiO2 sol-gel coatings studied by variable incidence infrared spectroscopy”, p. 220, Copyright (2002), with permission from Elsevier)

IS

Absorbance (a.u.)

37

1135

HS

900° C 700° C 650°C 500° C

350° C 1300 1200 1100 1000 900 1300 1200 1100 1000 900 Frequency (cm−1)

groups, coming from MTES in HS coatings, confirms the hybrid character of these films; this peak remains up to 650  C. Moreover, the Si–OH peak, at 960 cm–1, appeared only at 650  C for the HS films, coinciding with the complete elimination of the hydrophobic CH3 groups. This peak totally disappears for heat treatment temperatures above 700  C, for both types of coatings. For T > 700  C, therefore, the normal incidence spectra look very similar for both types of samples, with a welldefined peak at 1070 cm–1 and a smooth Sh feature on the high-frequency side, without detectable Si–OH or CH3 groups (Gallardo et al. 2002). In fact, Fig. 7 shows some evidence that the initial composition plays an important role on the final structure obtained after sintering. Especially in the case of HS films, the TO peak occurs at higher frequencies than in IS, and its area increases more pronouncedly with the sintering temperature, suggesting a higher polymerization and densification degree of the silica network. At heat treatment temperatures up to 500  C, sintering of the hybrid films is retarded by the presence of terminal methyl groups, but, for higher temperatures (>700  C), both spectra tend to become very similar: the Sh band suffers a decrease in intensity, indicating that porosity has been eliminated, and the organic peak at 1275 cm
1 disappears, indicating the loss of hybrid character of the HS films. However, oblique incidence spectra reveal additional important differences between IS and HS films, even after full densification, as seen in Fig. 8. In fact, at oblique incidence, the most important feature revealed is the activation of the LO-related vibrations, with the LO mode appearing as a strong peak at 1225–1260 cm
1 for the IS films, at all heat treatment temperatures, but only for

1136

IS

HS

900°C absorbance (a.u.)

Fig. 8 Evolution of FTIR spectra for coatings prepared from TEOS (IS) and MTES/ TEOS (HS), measured at 60 off-normal incidence, for different heat treatment temperature (Reprinted from Journal of Non-Crystalline Solids, 298, J. Gallardo, A. Durán, D. Di Martino, R.M. Almeida, “Structure of inorganic and hybrid SiO2 sol-gel coatings studied by variable incidence infrared spectroscopy”, p. 221, Copyright (2002), with permission from Elsevier)

R. M. Almeida and A. C. Marques

700°C 650°C

500°C 350°C

1400 1300 1200 1100 1000 1400 1300 1200 1100 1000 900 frequency (cm-1)

treatment temperatures above 700  C, in the case of HS films. Moreover, all the IS spectra and those of HS for treatment temperatures above 650  C show an unresolved band at 1200 cm
1, whereas the HS film spectra, for 350  C and 500  C, show a much stronger peak at 1140 cm–1 (Gallardo et al. 2002). The similarity between the spectra of HS, for 900  C and IS, for only 350  C, suggests that the structure of HS films, even after organic species removal and densification at 900  C, still retains some memory of the hybrid structure and has features distinct from IS films densified at 900  C. Evaporation from the liquid phase during film deposition is generally a fast process which triggers condensation reactions and at the very end the organization phenomena. A Si alkoxide bearing an epoxy functionality, namely, 3-glycidoxypropyltrimethoxysilane, has been employed in the preparation of hybrid thin films by EISA. Innocenzi et al. (2011) have studied the structural evolution of such hybrid sol (basic catalysis with NaOH) during the evaporation phenomenon by in situ time-resolved IR spectroscopy. The 950 cm-1 band of the epoxy group is generally used for assessing the presence of the epoxy functionality in organic resins. In hybrid SG materials, because of the high degree of overlapping in that region with the absorption band of silanols (Si–OH stretching), the band at 3050 cm
1 is more commonly used, but even in this case, overlapping with the C–H stretching region does not allow a very precise evaluation of the epoxy content. Innocenzi et al. (2011), using time-resolved FT-near-IR spectroscopy (range of 4600–4480 cm
1), showed that the epoxy groups do not react during the evaporation (the ring remains closed), while the silica structure shows only slight condensation and an increase in open cage-like species. The structure was found to evolve from a wet to a “soft-like” state.

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

1137

Multicomponent Silicates The preparation of multicomponent glasses by the SG process dates from several decades ago. SG is perhaps one of the best processing methods in order to obtain very homogeneous products. Nevertheless, the very different hydrolysis rates of the alkoxide precursors of the different oxides must be taken into account because they lead, if not appropriately controlled, to inhomogeneity of the resulting gel (Weinsenbach et al. 1991; Dawnay et al. 1995; Orignac and Almeida 1996; Guglielmi et al. 1998). IR spectroscopy has been a useful technique to verify the occurrence of phase separation, yielding a measure of the homogeneity (achieved through heterocondensation) versus phase separation (due to homocondensation), in the final glass network. FTIR studies have been performed on numerous multicomponent systems, in order to assess the relative amounts of specific types of chemical bonds and to determine the effects of various species and/or components on the glass properties and structure. Systems such as SiO2–TiO2 (Weinsenbach et al. 1991; Matsuda et al. 1993; Almeida et al. 1994; Kasgoz et al. 1994; Dawnay et al. 1995; Du and Almeida 1997; Guglielmi et al. 1998; Martins and Almeida 2000; Seco et al. 2000; Marques et al. 2003), SiO2–TiO2–AlO1.5 (Almeida et al. 1999; Yeatman et al. 2000), SiO2–TiO2–P2O5 (Orignac and Almeida 1996), SiO2–ZrO2 (Salvado et al. 1988; Okasaka et al. 1991; Colomban and Bruneton 1992; Saha and Pramanik 1993; Ricol et al. 1997; Neumayer and Carrier 2001), and SiO2–HfO2 (Neumayer and Carrier 2001), among several others, have been studied. The incorporation of additional components can significantly modify the final glass structure and the corresponding vibrational spectra. For instance, alkali additions to silicate glasses in the SiO2–ZrO2–Na2O system (Ricol et al. 1997) modify their spectral characteristics and the structure of the glass. Silicate glasses doped with rare earth and/or noble metal elements have also been the subject of several FTIR studies, e.g., SiO2–TiO2–ErO1.5–Ag (Almeida et al. 2002, 2004) and SiO2–TiO2–ErO1.5 plus SiO2–HfO2–ErO1.5 (Almeida et al. 2004).

Structural Evolution with the Heat Treatment, Temperature and Time Among the different multicomponent systems prepared by SG processing until now, the silica–titania-based materials are the most studied. The IR spectra of such films present a dominant band at about 1070 cm
1, due to the asymmetric stretching of the Si–O–Si bonding sequences and a weaker peak on its low-frequency side, at 950 cm
1, due to the overlap of Si–O– stretching vibrations in Si–O–Ti and Si–OH bonding environments (Almeida 1998). The Si–O–Ti stretching peak is characteristic of homogeneous gel/glass network regions, whereas the 1070 cm
1 peak is indicative of silica-rich regions in a phase-separated network (Almeida et al. 2002). Figure 9 shows the IR spectra of two of these films, before and after heat treatment at 600  C. After heat treatment, the major differences concern the

R. M. Almeida and A. C. Marques

Absorbance (a.u.)

1138

(b)

(a)

4000

3500

3000

2500 Wavenumber

2000

1500

1000

500

(cm–1)

Fig. 9 IR absorption spectra of 78.4% SiO2–19.6% TiO2–0.5% ErO1.5–1.5% Ag films: (a) as-deposited and (b) after heat treatment at 600  C, for 60 min (Reproduced with permission from Almeida et al. (2002))

elimination of the broad band due to SiO–H stretching of the silanol groups, at 3400 cm
1, and the decrease of the 950 cm
1 peak relatively to the 1070 cm
1 peak, plus a decrease of the shoulder at 1200 cm
1. Therefore, one may take the ratio of the Si–O–Ti to Si–O–Si peak intensities as a measure of the homogeneity (due to heterocondensation reactions), versus phase separation (due to homocondensation), in the film network. Figure 10 shows this evolution, as a function of heat treatment time at 600  C, for the same film: the continuous decrease in the degree of homogeneity is a clear indication of the progressive development of phase separation with time (Almeida et al. 2002).

Evolution of the TO Peak Frequency Figure 11 shows the evolution of the frequency of the TO peak, for silica and silica–titania SG films, with the heat treatment temperature. There is a minimum in the frequency values for both compositions at about 400–500  C, followed by a peak frequency increase for higher treatment temperatures (Almeida et al. 1994; Primeau et al. 1997; Innocenzi 2003). This frequency minimum may be associated with a maximum in the porosity of the films. Indeed, with the help of other characterization techniques (DTA, ellipsometry), it is possible to observe that the emptying of the pores of residual water and unreacted alkoxy groups, which is completed at T  450  C, results in a large volume fraction of porosity for these SG films. This porosity, in turn, induces a strain in the Si–O–Si bonds at the surface of the pores,

Characterization of Sol-Gel Materials by Infrared Spectroscopy

Fig. 10 Ratio of IR peak intensities as a function of heat treatment time, at 600  C, for 78.4% SiO2–19.6% TiO2–0.5% ErO1.5
1.5% Ag films (Reproduced with permission from Almeida et al. (2002))

1139

0.72 0.68 (Si-O-Ti)/(Si-O-Si)

37

0.64 0.60 0.56 0.52 0

5 10 15 20 25 30 35 40 45 50 55 60 Heating time (min)

1085 1082 1080 1080

Wavenumber / cm–1

1075 1078

1070 1076

1065

1074

1072

100S

1060

80S20T 1070 0

200

600 400 Temperature / C

800

1055 1000

Fig. 11 Position of main vibrational peak (from FTIR spectra) of a pure silica film and a film containing 30% titania (Reproduced with permission from Almeida et al (1994))

1140

R. M. Almeida and A. C. Marques

which leads to larger bridging angles and longer Si–O bonds, compared to bulk silica glass, causing a decrease of the TO peak frequency (Almeida et al. 1990a).

Comparison Between Different Systems Figure 12 shows the FTIR spectra, in the strong absorption region between 700 and 1400 cm
1, of several (Er3+-doped) silicate SG films, together with a pure silica film. The major difference, in this region, consists of a much weaker peak at 980 cm
1, for silica–hafnia films, assigned to Si–O–Hf asymmetric stretching (Neumayer and Carrier 2001), when compared to the Si–O–Ti peak at 940 cm
1, for silica–titania films (Almeida 1998). A mixed composition, silica–titania–hafnia, exhibits intermediate behavior regarding the intensity and frequency of this peak, which is characteristic of homogeneous gel/glass network regions. The position, intensity, and FWHM of the different peaks were determined through conventional peak-fitting procedures. The reason why the Si–O–Hf peak was much weaker than the Si–O–Ti peak might be due to a lower degree of homogeneity, in the former case, but it could

74.6% SiO2 - 10.0% TiO2 - 14.9% HfO2- 0.5% ErO1.5 79.6% SiO2 - 19.9% TiO2 - 0.5% ErO1.5 79.6% SiO2 - 19.9% HfO2- 0.5% ErO1.5 100% SiO2

Transmittance (a.u.)

945 cm-1

938 cm-1 973

1400

1300

1200

1100

cm-1

1000

900

800

700

Wavelength (cm-1)

Fig. 12 IR absorption spectra of different SG films: pure silica, silica–titania, silica–hafnia, and silica–titania–hafnia. All the monolayer films had similar thickness (300 nm) and were subjected to the same heat treatment (2 min at 900  C) (peak fittings are not shown for clarity)

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

1141

also arise in part from a lower IR activity of the Si–O–Hf vibrational mode, due to structural differences, such as different distributions of the Si–O–Hf and Si–O–Ti angles. The Sh/TO intensity ratio was also somewhat lower in the case of silica–hafnia films (0.19), compared to the case of silica–titania films (0.25), which might suggest that the former films densified more completely than the silica–titania ones (Almeida et al. 2004). No appreciable differences arose in these spectra due to doping with Er. It is also possible to observe that the dominant peak at 1070 cm
1, in the hafniacontaining sample (FWHM  84 cm
1), is substantially broader than the corresponding peak of the silica–titania film (FWHM  70 cm
1) or that of the silica–titania–hafnia (FWHM  66 cm
1) and also than that of pure silica films (FWHM  74 cm
1). This fact may indicate that the structure of the silica–hafnia glasses is more disordered, especially with respect to the Si–O–Si bond angle distribution, which must be broader (Almeida and Pantano 1990a). The incorporation of other elements, such as A1 or P (Fig. 13), also induces significant changes at the structural level: the Si–O–Ti bonds are replaced with Si–O–Al or Si–O–P upon A1 or P co-doping, respectively (Almeida et al. 1999). Thus, the peak near 940 cm
1, due to Si–O–Ti bonding sequences, shows a decrease in intensity, due to the decrease in the concentration of such bonds, when Al2O3 or P2O5 are added; the Si–O–Al and Si–O–P vibrational peak is probably hidden under the main Si–O–Si stretch at 1070 cm
1.

Fig. 13 IR absorption spectra of 80SiO2–20TiO2, 73SiO2-18TiO2–9AlO1.5, and 73SiO2–18TiO2–9PO2.5 SG films. All the films (consisting of four layers) had similar thickness (1 μm) and were subjected to the same heat treatment (2 min at 900  C, for each layer)

1142

R. M. Almeida and A. C. Marques

Influence of the Thickness and Number of Layers on the Microstructure of Multilayer Silicate Films This section contains a brief description of the microstructure evolution with the thickness and number of layers, for the same composition and heat treatment, in a multilayer silicate SG film. The number of layers which form a silicate SG film has been found to have a strong influence on the relative intensity of the TO and shoulder (Sh) peaks. Numerous experiments have led us to establish that, as the number of layers increases, the spectra shown in Fig. 14 exhibit a weaker absorption band at 1080 cm
1 (TO mode) and a slightly more intense band at 1180 cm
1; this is believed to be a mode of mixed longitudinal and transverse optical character (Hu 1980; Almeida and Pantano 1990a). The deconvolution of the spectra, for films consisting of 1 and 11 layers, revealed that the fractional area of the (TO + Sh) peaks is quite similar in both cases; this may suggest a correlation between the intensities of both peaks, in the sense that when the Sh peak increases (with the number of layers), the TO peak, at 1080 cm
1, tends to decrease, with the total (TO + Sh) peak area remaining approximately constant. Such an increase of the Sh peak can, in principle, be due to the fact that the overall film thickness increases with the number of layers, and for thicker films, the amount

1090 1088

Si-O-Si peak frequency

Absorbance (a.u.)

2 layers 3 layers 6 layers 11 layers

1086 1084 1082 1080 1078 1076 1074 0

2

4

6

8

10

12

number of layers

1400

1300

1200

1100 1000 Wavenumber (cm-1)

900

800

700

Fig. 14 IR absorption spectra (normalized to the Si–O–Ti peak) of 79.6% SiO2–19.9% TiO2–0.5% ErO1.5 multilayer films. Heat treatments consisted of 2 min at 900  C for each layer; plus a final treatment of 1 h at 900  C. The inset shows the evolution of the Si–O–Si peak frequency with the number of layers

37

Characterization of Sol-Gel Materials by Infrared Spectroscopy

1143

of residual pores, or other types of scattering centers, is higher than for thinner films. In fact, even if the residual pore concentration is independent of the film thickness (which increases with the number of layers), the total amount of pores crossed by the IR light becomes larger with increasing thickness. For thicker films, there will be a larger number of pores or other types of scattering centers in the light path, which will scatter the radiation in all directions, activating the LO mode and thus increasing the absorbance of the related Sh peak (Almeida and Pantano 1990a). For this reason, the observed increase of the Sh/TO intensity ratio in this case does not necessarily indicate larger pore concentration or larger volume fraction of porosity, Vp. Another possible source of such Sh/TO intensity ratio increase might be connected with the scattering caused by layer boundaries and striations, which play a more significant role in multilayer films, causing LO mode activation, as well. Also, the frequency of the TO peak shows a significant shift to higher values with increasing number of layers (inset of Fig. 14), which is difficult to explain. Although, for an increasing number of layers, each previously deposited layer is treated for 2 min at 900  C an increasing number of times, the final heat treatment of the multilayer film for 1 h at 900  C should even out these effects. A longer heat treatment would in principle induce a higher degree of densification and a simultaneous stiffening of the Si–O–Si bonds, consequently increasing the TO peak frequency values. The study of multilayer films, subjected to the same heat treatment and consisting of equal number of layers, but with different thicknesses (prepared from sols with different EtOH/precursor ratio, RE), reveals quite similar results when compared to the effect of different thicknesses caused by a different number of layers. As shown in Fig. 15, for the single-layer films only, the thicker film has also a higher TO peak

Fig. 15 IR absorption spectra (normalized to the Si–O–Ti peak) of 79.6% SiO2–19.9% TiO2–0.5% ErO1.5 single-layer films. Both films are similar, except in thickness. The thickness difference arises from the different RE values used in the sol preparations

1144

R. M. Almeida and A. C. Marques

frequency (1066 cm
1, for the thicker film, and 1063 cm
1, for the thinner one) and higher Sh/TO intensity ratio (0.25, for the thicker film, and 0.18, for the thinner one). The shift of the TO peak toward slightly higher frequency values, in the thicker film, could be explained, in this case, by the fact that a larger cross-linking degree is caused by the lower RE value used on the preparation of the sol, leading to a stiffening of the Si–O–Si bonds. The larger Sh/TO peak intensity ratio could be explained by the number of scattering centers in the light path, similarly to the previously used argument.

Nonsilicates and Non-oxides Silicates are the most common of all glasses, but there are many other types of oxide, as well as non-oxide glasses, including, most notably, halides and chalcogenides. Nonsilicate oxides, such as zirconia (ZrO2), titania (TiO2), or alumina (Al2O3), processed by SG, have been widely studied due to their promising properties, like chemical durability, alkali resistance, high refractive index, photocatalytic activity, bioactivity, thermal stability, mechanical strength, refractoriness, or optical properties (Izumi et al. 1989; Atik et al. 1995; Jokinen et al. 1998; Djaoued et al. 2002a). These properties strongly depend on the microstructure, which can be studied by IR absorption spectroscopy. Zirconium alkoxides, such as Zr isopropoxide and Zr n-butoxide (Izumi et al. 1989; Neumayer and Carrier 2001) or Zr acetylacetonate (Salvado et al. 1988), Zr tetraoctylate (Izumi et al. 1989), and Zr formate (Saha and Pramanik 1993), are just some of the precursors that can be used for the synthesis of ZrO2 gels. The choice of the starting material has a large influence on the stability of the sols and, consequently, on the final SG products. IR spectroscopy is useful in the detection of organic compounds (C–H groups), OH groups and crystalline phases, present in this kind of materials. Before heat treatment, the IR spectra show absorption bands due to Zr–O–C bonds, near 1450 and 1560 cm
1, an absorption peak near 2900 cm
1, assigned to C–H bonds and weak bands due to Zr–O–Zr species, located near 750 and 500 cm
1. The evolution of the spectra with the thermal treatment of the samples shows that the absorption due to Zr–O–C and C–H disappears, while that of Zr–O–Zr increases and becomes sharp, with the time and temperature of heat treatment (Izumi et al. 1989). Neumayer and Carrier (2001), on the other hand, reported that, after a heat treatment at 400  C for 1 h, the IR spectra showed a broad absorption band from 600 to 200 cm
1, peaking at 378 cm
1, which was assigned to Zr–O vibrations. The width of this band suggested a significant dispersion of vibrational states, characteristic of an amorphous network. Small crystallites of tetragonal (t) ZrO2 and monoclinic (m) ZrO2 may be nucleated from the amorphous gel matrix during heat treatment. Indeed, for higher heat treatment temperatures, peaks at 436 and 158 cm
1, typical of ZrO2 (t), are present in the spectra. For temperatures above 800  C, the transformation from ZrO2 (t) into

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ZrO2 (m) is clearly visible by the progressive appearance of the absorption bands at 722, 574, 490, 409, 343, 258, and 228 cm
1, characteristic of ZrO2 (m) (Neumayer and Carrier 2001). Numerous titanium alkoxides are used in SG chemistry, in order to prepare inorganic titanium oxides. They are, for instance, Ti tetra-iso-propoxide (TIPT) and Ti tetra-n-butoxide (TBOT). The SG transformation mechanisms of TIPTderived coatings have been qualitatively studied by FTIR spectroscopy (Burgos and Langlet 1999), allowing a good knowledge of the hydrolysis and condensation reactions of the alkoxide. The IR spectra of crystalline TiO2 have been widely reported and consist of the absorption peaks at 433–438 and/or 489–496 cm
1, depending on the temperature of the heat treatment and, therefore, on the TiO2 crystalline phase present, anatase and/or rutile. Moreover, before heat treatment, there are still ethanol (ethyl groups) and isopropyl groups remaining, represented by the peaks at 3222 cm
1 (vOH), 1064 cm
1 (vC–O), 2906 cm–1 and 2855 cm–1 vCH2 , CH3 Þ , and 1355 cm–1 plus 1448 cm–1 δCH2 , CH3 (Djaoued et al. 2002b), in addition to a broad band between 400 and 800 cm
1, which corresponds to the formation of Ti–O and Ti–O–Ti bonds with C and/or hydroxyl groups incorporated (Alam and Cameron 2002). All these bands disappear for T > 300  C, except the broad band at 400–800 cm
1, whose intensity increases with the temperature of the heat treatment performed (Alam and Cameron 2002). After a heat treatment at 500  C, there is the formation of a well-defined band around 434 cm
1, assigned to the Ti–O–Ti stretching vibration in the anatase phase, which tends to increase as the temperature increases up to 600  C. For T > 700  C, the band at 434 cm
1 decreases and eventually disappears, for T > 800  C. In the meantime, a new band at 485 cm
1, ascribed to Ti–O–Ti stretching vibrations in the rutile phase, appears (Alam and Cameron 2002; Djaoued et al. 2002b). Despite the fact that all the studied films start to crystallize at 400  C, Djaoued et al. (2002b) found a significant influence of complexing agents, present in the coating solution, on the phase transition temperature. For instance, acetylacetone (AcAc) and its mixture with acetic acid were found to stabilize the anatase phase, even up to temperatures as high as 1000  C. Furthermore, Djaoued et al. (2002a) reported a low-temperature method (T  100  C) which, by removal of OH groups, allows a rearrangement of the Ti–O network and promotes the crystallization of titania at very low temperature. This method consists of the deposition of the films, followed by a hot-water treatment for selected times. By using IR absorption spectroscopy, it was possible to study the evolution of the bands assigned to vOH, δOH, Ti–OH and Ti–O, before and after the hot-water treatment. They showed that this treatment promotes a decrease of the intensity of the vOH band and the early appearance of the Ti–O band around 453 cm
1, even for temperatures as low as 100  C. IR absorption spectroscopy has also been used with the aim of studying the in vitro bioactivity of SG-derived TiO2 films deposited by dip coating (Jokinen et al. 1998). The IR spectra of the TiO2 films showed an increase of the absorption

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peaks related to vibrational modes of the carbonate ion (1650–1300 cm
1, 873 cm
1) and phosphate ion (962 cm
1, 556–597 cm
1, 1096 cm
1, 1112 cm–1), after their immersion in a simulated body fluid (SBF). This indicates the existence of calcium phosphate phases and, therefore the formation of bone-like hydroxyapatite, with the immersion of the TiO2 films in the SBF, evidencing their in vitro bioactivity. The processing of mesoporous TiO2 films by EISA has been studied by (in situ) time-resolved IR spectroscopy in order to follow two of the main stages of selfassembly: evaporation and post-deposition drying. The nanostructured TiO2 cast films have been prepared by using a triblock copolymer (Pluronic F127) and TiCl4 as the titania precursor, subjected to a heat treatment up to 165  C. The evaporation phenomena can be directly observed by FTIR because there is no overlap of the IR signals from the Ti precursor and surfactant. After drying at 165  C, there was evidence from the IR spectra for the complete removal of water and for triblock copolymer crystallization as soon as the water evaporates, which was evidenced by transformation of the typical broad band within the range of 1200–1000 cm
1 into the triplet signature of the Pluronic F127 crystalline phase (Fig. 16). Crystalline micelles are therefore detected within the TiO2 mesopores. Such crystallization was found to be reversible when cooling to room temperature (Innocenzi et al. 2010). Alumina is a well-known high-temperature insulating material, and therefore, it is of value the use of FTIR in the study of the sequence of transformations which occur during heat treatment of the gel material, at different temperatures. On the other hand, the low-temperature phases of alumina are important in catalysis, due to their high specific surface areas and the large number of defects in their crystalline

Fig. 16 Time-resolved IR spectra of a titania mesostructured film as a function of drying temperature. The spectra have been recorded during in situ measurements from 25 up to 165  C, with steps of 10  C (Reprinted with permission from Innocenzi P., Malfatti L., Kidchob T., Grosso D. Controlling the processing of mesoporous titania films by in-situ FTIR spectroscopy: getting crystalline micelles into the mesopores, J. Phys. Chem. C 2010; 114: 10806–10811. Copyright 2010 American Chemical Society)

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structure (Wang et al. 1999). There are several Al2O3 crystalline phases (η, γ, δ, κ, θ, and α), but since the transition temperature range between them is sharp, most of them appear always mixed. The transformation between the different Al2O3 phases strongly depends on the precursors and the heat treatment performed for their stabilization. Wang et al. reported the study of the low-temperature phases of alumina, in particular the evolution of hydroxyl species with temperature, measured by FTIR spectroscopy. They found that, for the fresh gel sample dried at 70  C, there was an intense and wide absorption band at 3455 cm
1, assigned to OH groups on the γ-AlO(OH) surface. With the increase of the heat treatment temperature up to 800  C, this band tends to decrease in intensity, due to the removal of water molecules and the boehmite (AlO(OH)) transformation into γ-Al2O3. At T = 800  C, OH bands still remain in the spectra, which shows the high capacity of the low-temperature alumina phase for retaining hydroxyls in its structure, indicating a high specific surface area, an important property in catalysis. O–Al–O bonds also contribute significantly to the IR spectra. Stangar et al. (2002) reported the IR absorption spectra of alumina xerogel powders (in KBr pellets), heat treated at temperatures ranging from 450–550  C, which exhibited two bands, at 640 and 850 cm
1, attributed to vibrational modes of O–Al–O groups. At 600  C, the splitting of these bands decreased (670 and 780 cm
1), and both bands tend to degenerate into a single broad band at 710 cm
1, when the heat treatment temperature was above 650  C. SG-derived ceria–titania films are of importance due to their application as antireflective coatings and electrochromic layers and in self-cleaning glasses. In particular, thermal treatment of ceria–titania films in air at high temperature up to 700  C could lead to the formation of cerium titanate phases such as CeTi2O6. The combination of far IR spectroscopy and X-ray diffraction yields a powerful tool to investigate the crystallization process and identify the phases formed as a function of temperature and composition (Kidchob et al. 2009). However, it should be stressed that a far IR absorption band peaking around 250 cm
1 starts to be observed after annealing at 300  C, increasing its intensity for higher temperatures, which suggests that small crystallites of ceria and cerium titanate, which are not yet detectable by X-ray diffraction, have already formed at these lower temperatures, revealing the sensitivity of the FTIR technique for detecting early stage crystallization phenomena. Information on SG processing and characterization of non-oxides is very scarce. An example of reported IR absorption spectra of non-oxide SG materials is for sulfides, namely, GeS2, with characteristic Ge–S IR peaks at 375 and 430 cm
1, assigned to asymmetric stretching vibrations of GeS4 and S3Ge–S–GeS3 units (Frumarová et al. 1999). The synthesis of GeS2 without contamination by oxide (GeO2) is not trivial. Middle and far IR absorption spectra were reported and discussed by Martins et al. (1999), in order to detect the presence of GeO2 in GeS2-based films, after using different Ge precursors, such as germanium tetrachloride (GeCl4) and germanium tetraethoxide and different sources of S, such as thiourea, thioacetamide, and hydrogen sulfide (H2S). For films prepared by the reaction of thioacetamide and germanium tetrachloride, dissolved in ethanol, the

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absorption bands of Ge–O bonds at 870 and 560 cm
1 were not detected, and therefore, only Ge–S bonds were found in the IR spectrum (Martins et al. 1999). Xu and Almeida (2000) reported a successful preparation of GeSx SG films (with a content of GeO2 of only 6%), using GeCl4 and H2S as the main precursors and toluene as the solvent. In this work, the concentration of Ge–O bonds in the sulfide films was estimated based on Beer’s law: A ¼ ebc

(10)

where A is the absorbance, e is the molar absorptivity, b is the film thickness, and c is the concentration of absorbing species, in this case Ge–O bonds. The values found, for the concentration of Ge–O bonds, were 100%, for a GeO2 film and 6%, for a GeSx film.

Conclusions The IR spectra of SG materials contain considerable information about their composition, structure, and properties. Therefore, IR spectroscopy has been widely applied for silicate materials, but also for other oxides, as well as some non-oxide materials. Some examples of the rich information provided by IR spectroscopy, described in this chapter, include the structural evolution during the different stages of SG material processing, including complex phenomena occurring during EISA, the semiquantitative detection of OH species and porosity, the analysis of structural homogeneity, phase separation and densification degree, and the study of structural differences between thin films, thicker multilayer films, and bulk SG samples, as well as between SG-derived purely inorganic and hybrid materials. In addition, in situ and time-resolved IR spectroscopy analysis enhances the ability to follow the reactions and physical phenomena occurring at the various stages of SG processing.

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Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies

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Maurizio Montagna

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Inelastic Light Scattering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Raman and Brillouin Instrumentation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Vibrational Dynamics of Aerogels: Fractons . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Densification of Silica Xerogels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Characterization by Waveguided Brillouin Spectroscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Raman Spectroscopy of Glass Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Raman Spectroscopy of Nanocrystals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1154 1154 1158 1158 1166 1170 1173 1175 1179 1179

Abstract

The use of Raman and Brillouin spectroscopies in sol-gel-derived materials is reviewed. It covers a quite vast domain of investigation, from the basic glass science to the characterization of materials produced for many different applications. The theory of inelastic light scattering is briefly presented and the basic physical mechanisms involved in Raman and Brillouin spectroscopies are discussed. The instrumentation for measurements by visible, ultraviolet, and X-ray excitation is described. The vibrational dynamics in low-density aerogel is discussed in terms of a random fractal model, where the low-frequency acoustic vibrations are distorted phonon-like extended propagating modes, whereas highfrequency modes can be spatially localized. As a function of the vibrational frequency, the phonon wavelength and mean free path depend on the size of the porosity. The dependence of the sound velocity and attenuation on the densities is M. Montagna (*) Disordered System Laboratory, Department of Physics, Università degli Studi di Trento, Trento, Italy e-mail: [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_34

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measured by Brillouin spectroscopy. Attenuation in low-density sol-gel-derived solids is nearly temperature independent and is due to structural disorder instead than to anharmonicity, as it is in crystals and compact glasses. It is shown that Raman and Brillouin spectroscopy can be used to follow the different steps of densifications from the wet-gel to the compact glass during thermal annealing. Waveguided spectroscopies for dip- and spin-coated films are described. Low-frequency Raman spectra from the localized acoustic vibration in nanocrystals are presented in sol-gel-derived glass ceramics.

Introduction Raman spectroscopy is a powerful characterization technique in materials science, because the vibrational dynamics gives rich information on the molecular structure. Furthermore, it is a nondestructive technique of low cost and simple analysis of the data. It is widely used for the study of inorganic, organic, and hybrid materials processed by sol-gel technology in all the steps of the process: from the starting solution through hydrolysis and polycondensation, gelling, drying of water and solvents, and high-temperature annealing to form glass and ceramics. About 100 papers per year on studies of sol-gel materials by Raman scattering are now appearing. Bulk samples but also films are characterized by the Raman technique. In the latter case, the exciting laser light is usually coupled to the waveguide by butt coupling, prism coupling, or writing a grating in the film. Brillouin scattering, even if less extensively used as a characterizing technique, gives important structural information on materials produced by sol-gel, since the sound velocity strongly depends on the porosity. By prism coupling, Brillouin scattering in planar waveguides can be measured. Inelastic light scattering has been also extensively used to study the vibrational properties of fractal systems: silica aerogels, with their porosity extending on a wide range of sizes, are very nice examples of fractal systems. The peculiar vibrational dynamics of these systems, with spatially localized vibrational modes, is at the origin of the extremely low-measured thermal conductivity. In this chapter, the application of inelastic light scattering to the study of the vibrational dynamics of sol-gel materials in bulk and planar waveguide forms will be described.

Inelastic Light Scattering Raman and Brillouin scattering are two photon processes. The electromagnetic field exchanges an energy ℏω = hυin
hυout and a wavevector q = kin
kout with the system under study. υin and υout are the frequencies of the incoming and scattered photons, respectively. kin and kout are the wavevectors of the incoming and outgoing

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Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies

1155

photon, respectively, with k = 2π/λ = 2πn/λ0, where n is the refractive index and λ and λ0 are the wavelengths in the scattering system and in vacuum, respectively. The units which form the system, atoms, ions, and bonds are polarized by the exciting field and by the fields produced by the other dipoles, through a dipole-induced dipole mechanism (DID). These units irradiate fields at the exciting frequency (elastic or Rayleigh scattering), but also inelastic scattering is produced by the thermal fluctuations of the system. The scattered intensity is proportional to the fourth power of the frequency and to the square of the sum of the electric fields irradiated by all the polarized units. In a quantum mechanical approach to the polarizability of the basic units, the inelastic scattering process is the destruction of a photon, due to an electronic transition to an excited level and the creation of a photon of different energy, assisted by the creation or destruction of one or more vibrational quanta. If the system is transparent to the exciting and scattered light, a virtual transition to the intermediate electronic level occurs (normal Raman scattering). On the contrary, when the photon energy is resonant or near resonance with an electronic transition, one has resonant Raman scattering. The intensity in the various scattering configurations is given by the Fourier transform of the correlation functions (Benassi et al. 1993): Gαβγδ ðtÞ ¼

XD   E Πiαβ ðtÞΠjγδ ð0Þexp iq  ri ðtÞ
rj ð0Þ ;

(1)

ij

where α, β, γ, and δ are Cartesian indices and Πiαβ are the instantaneous atomic polarizability tensor components of the ith atom, at time t and position ri(t). Equation 1 accounts for the elastic (Rayleigh) and inelastic (Brillouin and Raman) light scattering. After expansion in the atomic displacements, ui(t), from the equilibrium positions, Ri, ri ðtÞ ¼ Ri þ ui ðtÞ;

(2)

the Brillouin and one-phonon Raman contributions are given by the linear terms in the displacements u (Benassi et al. 1993): X i

   Πiαβ ðtÞeiqri ðtÞ  

¼ l ph

XX i

e

iqRi

μ



iqμ Πiαβ

 eq

þ



Qiαβ, μ



eq

uiμ ðtÞ;

(3)

where α and β are the direction of polarization of the incoming and outgoing photons, respectively, and (Πiαβ)eq is the polarizability of the ith atom, for all atoms at the equilibrium position, and   X @Πiαβ ¼ Qiαβ, μ eq @uiμ l

! exp½iq  ðRl
Ri Þ: eq

(4)

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M. Montagna

The first term in the sum of Eq. 3 describes the polarizability modulation on the system by a vibrational mode that displaces the masses from the equilibrium positions. The second term in the sum in Eq. 3 accounts for the polarizability modulation of the ith atom caused by the displacements of the surrounding atoms. For molecules with a small size, a, (qa  1), the fields scattered by all the polarizable units are in phase. In this case a single quantity, the polarizability Пαβ of the molecule determines the Raman spectra. The symmetry of the molecule and of the Пαβ tensor is fundamental for the calculation of the symmetry of the vibrational modes and of the intensity of the relative Raman and IR absorption bands (Long 1977). In harmonic crystals, the vibrations are phonons, i.e., plane waves involving the motion of all atoms in the sample. Phonons are described by a frequency ωph and a wavevector kph. The light is scattered by the whole illuminated volume seen by the detector, and fields irradiated by all polarizable units interfere. As a consequence, one-phonon scattering processes (first-order Raman scattering) are subjected to the selection rules: ωph ¼ 2π ðυin
υout Þ,

kph ¼ q;

(5)

where the plus and minus signs refer to Stokes and anti-Stokes scattering, respectively. In light scattering experiments, the exchanged q is much smaller than phonon wavevectors at the boundary of the Brillouin zone. Since the dispersion curves of optical phonons are flat at the zone center, Raman scattering is q independent, and q = 0 can be taken in Eq. 3. Therefore the analysis of the Raman scattering by optical phonon in crystals is based on the structure and symmetry of the unit cell and on the symmetry of its Παβ, in a way very similar to that of molecules. On the contrary, Brillouin scattering of acoustic phonons is q dependent. By combining the two above selection rules, one obtains ωLph, T ¼ 2kin υL, T sin ðϑ=2Þ;

(6)

where ϑ is the scattering angle, within the sample, and υL and υT are the longitudinal and the (two different, in general) transverse sound velocities, having noted that kin ffi kout. In harmonic glasses, the vibrational normal modes are no longer phonons, due to disorder and lack of the translational symmetry of crystals. However, the Brillouin spectra of glasses are very similar to those of crystals, with sharp peaks, indicating that the low-frequency acoustical modes are nearly plane waves. The Brillouin peaks are found at about 10–40 GHz with a line width, which is a measure of the phonon mean free path, of the order of 100 MHz in simple glasses at room temperature. The temperature-dependent dynamical effects, as anharmonicity and glass relaxation, are comparable to the structural effects, which produce scattering of the plane acoustic waves by the disordered structure. In X-ray inelastic scattering experiments, where the Brillouin peak is in the THz range, the line width is an important fraction of the

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Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies

1157

peak frequency. Furthermore, it is nearly temperature independent, indicating that acoustic waves in the THz range are strongly scattered by the glass-disordered structure. Aerogels and most sol-gel-derived materials, having a residual porosity, are strong scatterers of the acoustic modes. The Brillouin line widths in these systems are larger than in the homogeneous bulk glasses and are nearly temperature independent. This subject is discussed in “Vibrational Dynamics of Aerogels: Fractons.” The Raman selection rule on the wavevectors in Eq. 5, valid for crystalline systems, does not hold for glasses, since the vibrational modes of glasses are different from planar waves. Instead of the sharp peaks of crystals, due to the optical phonons, broad bands appear in the Raman spectra of glasses. All vibrational modes are active, with different extents, in Raman scattering. Therefore, it is common to write the Raman intensity in VV-polarized (α = β) and HV-polarized (α 6¼ β) as (Shuker and Gammon 1970) I αβ ðωÞ ¼

nðω, T Þ þ 1 Cαβ ðωÞρðωÞ; ω

(7)

where n(ω, T ) is the Bose–Einstein population factor, C(ω) is the Raman coupling coefficient, and ρ(ω) is the density of vibrational states. For the high-frequency modes, the analogous of the optical modes of the crystals, the Raman activity is mainly related to the symmetry, as for molecules. Also in glasses, it is possible to deduce the symmetry of optical modes from the activity in Raman and IR absorption spectra. In particular, an important quantity is the depolarization ratio, i.e., the ratio of intensity of the HV- and VV-polarized spectra. In any case, there are not strong selection rules, due to disorder. The activity of the low-frequency modes, of acoustic-like nature or mixed acoustical and optical nature, has a more subtle origin, since it is intrinsically related to the presence of electrical and mechanical disorder (Martin and Brenig 1974). It is the disordered structure that does not allow a complete destructive interference of the scattered fields, as it occurs in crystals, where the acoustical phonons do not contribute to the Raman scattering (Benassi et al. 1995). The electrical disorder is caused by the disordered space distribution of the polarizability, as in the case of heavy ions in a silicate glass (Benassi et al. 1991). Mechanical disorder is the deviation of the vibrational mode patterns from the plane wave shape of phonons (Martin and Brenig 1974). In particular, a depolarized broad peak is present in the Raman spectra of all glasses, the boson peak, at frequencies in the range 20–60 cm
1. It corresponds to an excess in the density of states with respect to the ω2 Debye value, as measured in inelastic neutron scattering (INS) experiments. The nature of the vibrational modes at the boson peak frequencies has been extensively debated and many different models have been proposed (Fontana and Viliani 1998). Finally, of particular interest for the study of sol-gel-derived materials is the low-frequency Raman scattering of the acoustic vibrations of nanoparticles embedded in a glass matrix. This subject is discussed in “Raman Spectroscopy of Nanocrystals.”

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Raman and Brillouin Instrumentation The evolution of Raman instrumentation has been recently reviewed (Lewis and Edwards Howell 2001), from the first measurements by C. V. Raman, when the spectra were excited with a mercury lamp and recorded with a small prism spectroscope equipped with a photographic plate (Raman 1928), to the new high-resolution microscopes, in Raman scanning near-field optical microscopy (Adar 2001). This review includes Raman microscopy (Baldwin et al. 2001), Raman imaging (Treado Patrick and Nelson Matthew 2001), the adaptation of Raman spectrometry to industrial environment (Slater Joseph et al. 2001), Raman spectroscopy of catalysts (Wachs 2001), and process of Raman spectroscopy (Lewis 2001). Therefore, we restrict the discussion here to Brillouin equipments. Brillouin spectrometry requires setups with high resolution, contrast, and luminosity, since Brillouin lines are very close to the laser excitation, with typical shifts of 10–30 GHz, and very weak, about 10
10–10
12 times weaker than the laser line. On the other hand, a spectral range limited to a few GHz is required. Therefore, Brillouin spectrometers are usually multipass Fabry–Perot (FP) interferometers. Multipass tandem interferometers, based on two-plane FP (Sandercock 1970, 1976, 1978) or a plane and a confocal FP (Sussner and Vacher 1979; Vacher et al. 1980) in sequence, are used to fulfill the request of high contrast and resolution together with a sufficient spectral range. A light-modulation technique at microwave frequencies by a LiNbO3 crystal is usually employed for accurate measurements (Δυ/υ  10
4) of Brillouin line shifts (Sussner and Vacher 1979). A double spectrometer with a focal length of 2 m, and a resolution of about 0.015 cm
1 at 514.5 nm in double pass, is also used. It operates at the 11th order of high-quality ruled gratings with 300 lines/mm. It allows one to measure together Brillouin and low-frequency Raman spectra (Mazzacurati et al. 1988). A new spectrometer for high-resolution and high-contrast scattering spectroscopy in the ultraviolet has been constructed in L’Aquila, Italy. The instrument has two coupled 4 m-focal grating monochromators, with large-dimension gratings (204 408 mm2) ruled with 31.6 grooves/mm, working at the 230th order at λ = 266 nm and at the 115th order at λ = 532 nm, with a spectral range of 160 cm
1 (Caponi et al. 2004). Unfortunately, the L’Aquila earthquake in 2009 damaged it seriously. Brillouin spectroscopy in the THz frequency range is now possible by inelastic X-ray scattering at the very-high-energy-resolution IXS beamlines at the European Synchrotron Radiation Facility, Grenoble, France (Ruocco and Sette 2001; Scopigno et al. 2002).

Vibrational Dynamics of Aerogels: Fractons Silica aerogels have been considered as a model system for the study of the vibrational dynamics of random fractals. This, in turn, provided a deep knowledge of the structure of silica aerogels in terms of the preparation process. Fractals are selfsimilar systems described by a fractional space dimension D (Mandelbrot 1982).

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Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies

1159

Scaling is the basic concept for dealing with disordered media having self-similar symmetry (Alexander and Orbach 1982; Rammal and Toulouse 1983). The mass density of a fractal scale is M(R) / RD. This means that, starting from an arbitrary point on the fractal, the mass included in a sphere of radius R around the point increases as RD, instead of as Rd, d being the Euclidean dimension (d = 3 for the three-dimensional space). The structure of silica aerogels, with porosity on a wide range of distances, is well described by a fractal model (Vacher et al. 1988). The system is built by elementary grains of dimension a, which form a fractal structure that becomes statistically homogeneous at a crossover length, ξco, of the order of the maximum pore size. Small-angle X-ray scattering (SAXS) (Schaefer and Keefer 1986) and neutron scattering (SANS) (Vacher et al. 1989) have revealed welldeveloped fractal behavior over several orders of magnitude in length. In particular, samples with identical basic structure have been prepared with different densities and thus different crossover lengths, ξco (Vacher et al. 1988). In these systems, the acoustic vibrations are distorted phonon-like propagating modes for wavelengths λ larger than ξco and frequencies smaller than a first crossover frequency ωco1. In fact, the disordered porous structure is statistically homogeneous on these long-range scales. As in compact glasses, these phonon-like modes can be characterized by a wavevector q(q = 2π/λ), and a linear relation, ω = υL,T q, is observed between frequency and wavevector. The transverse, υT, and longitudinal, υL, sound velocities depend on the mass and bond distribution in the structure. They are smaller than those of bulk silica and decrease as the density of the aerogel decreases. The long-wavelength sound velocities in the phonon regime scale with ~ the density as υL, T / ρðD=d
1Þ=ðd
DÞ (Alexander et al. 1993). For wavelengths smaller than ξco and for the corresponding frequencies higher than ωco1, acoustic vibrations are fractons, i.e., highly disordered modes, localized on a volume of the order of λ3. It is still possible to define a (mean) wavelength λ, but a wide q distribution is present in the spatial Fourier transform of the modes (Montagna et al. 1990; Mazzacurati et al. 1992). Other two lengths useful for describing fractons are (i) the localization length LL, which describes the exponential or super-exponential decay of the envelope of the squared displacements from the center of localization of the mode (Petri and Pietronero 1992), and (ii) the scattering length LS, defined as the mean free path of a planar wave of frequency ω. LS can be obtained by measuring the width in q of the dynamical structure factor S(q, ω) in an inelastic scattering experiment, such as optical, X-ray, or neutron Brillouin scattering. In fractals, the three lengths λ, LL, LS follow the same scaling law (Alexander and Orbach 1982; Aharony et al. 1987): λ LL LS / ω
ðd =DÞ : ~

(8) ~

The fracton density of states scales with frequency as GðωÞ  ωd
1 ; where d~ is the fracton or spectral dimension. It describes the dynamics of the fractal and is different from the fractal dimension D. ωco1 and ξco depend on the density and are
ðd~=DÞ related by the relationship ξco / ωco1 :

1160 8000 Intensity (Counts/0.5 sec)

Fig. 1 Backscattering Brillouin spectra for six silica aerogel densities (in kg/m3). The relative intensities, not adjusted for sample turbidity, are otherwise significant. IW is the full instrumental width at half-height. The central portion of the spectra, affected by the elastic line, was removed for clarity (Reprinted figure with permission from Courtens et al. (1987); copyright (1987) by the American Physical Society)

M. Montagna

Q = 180°

407

6000 360

1W

103

4000

330 186

228

2000

0

0

1

2 3 Frequency w/2p (GHz)

4

5

The nature of the vibrational modes of aerogels in the frequency range around ωco1 was well assessed by Brillouin scattering measurements (Courtens et al. 1987, 1988). By varying the sample density and the exchanged q vector, Brillouin scattering covers a quite broad range of acoustic behavior, from disordered phonon-like modes with qξco  1 to fractons with qξco > 1. Figure 1 shows the backscattering Brillouin spectra of samples with different densities. Figure 2 shows the spectra with different exchanged q values for a sample with a density ρ = 186 kg/m3. The samples are part of a series of unoxidized neutrally reacted gels characterized by SANS (Vacher et al. 1988). They were obtained by hydrolysis of tetramethyl orthosilicate (TMOS) dissolved in water–methanol mixtures under initially neutral conditions. After gelling and long aging, the material was hypercritically dried. The spectra, excited with the 514.5 nm argon-laser line, were measured with a six-pass tandem interferometer of about 3 cm spacing (Sandercock 1978). Figure 1 shows that the longitudinal phonon velocity υL(ρ) decreases with decreasing ρ (increasing ξco). Note that the frequency shift of the peak is not completely due to the change of the sound velocity. In fact, also the refractive index changes with the density. Its value is well approximated by an interpolation between the refractive index of bulk silica and that of vacuum: nðρÞ
1 ¼ ðnSiO2
1Þ

ρ ρSiO2

;

(9)

with nSiO2 ¼ 1:46 at 514.5 nm and ρSiO2 ¼ 2200 kg=m3 : With decreasing density, υL decreases, but the attenuation, measured by the line width of the Brillouin peak, increases. The high-density samples show a phonon-like spectrum with a line width Γ smaller than the peak frequency υP. The crossover to fracton dynamics appears as a strong increase of the attenuation. For the sample with ρ = 186 kg/m3, the fracton regime has been almost completely reached, since Γ ffi υP. The presence of a phonon–fracton crossover in the dynamics of aerogels appears even more evident in Fig. 2. At small scattering angles, Brillouin scattering from acoustic phonon is observed. As the scattering angle increases, the exchanged q and the phonon

Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies

Fig. 2 Brillouin spectra for a silica aerogel with ρ = 186 kg/ m3 at four internal scattering angles θ, corresponding to k/2π ranging from 17,900 to 40,300 cm
1, for increasing θ (Reprinted figure with permission from Courtens et al. (1987); copyright (1987) by the American Physical Society)

1161

Q = 180°

Intensity (Arbitrary Units)

38

Q = 91°

Q = 67°

Q = 53° 0

3 1 2 Frequency w /2p (GHz)

4

frequency increase, approaching the crossover to the fracton regime. A detailed analysis allows one to obtain the dependence of the sound velocity, the attenuation, and the crossover frequency, on the sample density. The density of states of aerogels has been studied by low-frequency Raman spectroscopy (Boukenter et al. 1986; Tsujimi et al. 1988). As shown in Eq. 7, the Raman intensity depends on a coupling coefficient C(ω), which measures the Raman activity of each mode, preventing the possibility of a direct determination of the density of states. For fractal systems, the coupling coefficient is expected to have a scaling law CðωÞ  ωx ; since all physical quantities should scale in a self-similar system. In fact, the fractal model is confirmed by the Raman spectra (Boukenter et al. 1986; Tsujimi et al. 1988), which show a power law for the frequency dependence of the intensity. Figure 3 reports the Raman spectra of four neutrally reacted oxidized aerogels with densities between 158 and 357 kg/m3. The spectra are depolarized, i.e., the polarizations of the exciting and scattering light are crossed. The measurements were performed with a six-pass tandem Sandercock Fabry–Perot interferometer (Sandercock 1978), by exciting with the 514.5 nm line of an Ar+ laser, operating on a single mode at 250 mW. The data were taken by using different mirror

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RAMAN SUSCEPTIBILITIES (ARB. UNITS)

NEUTRALLY REACTED OXIDIZED

10

−0.39

−0.35

158

−0.36

−0.36

201

260

357 1 1

10 FREQUENCY (cm−1)

Fig. 3 The Raman susceptibilities I(ω)/n(ω) for four silica aerogel samples designated by their densities in kilograms per cubic meter. The corresponding acoustic correlation lengths ξac are 750, 480, 300, and 170 Å, in order of increasing densities (Courtens et al. 1988). The straight lines are fits with the indicated slopes, while the thin curves are guides to the eye. The different symbols correspond to the four different spacings L of the FP interferometer (Reprinted figure with permission from Tsujimi et al. (1988); copyright (1988) by the American Physical Society)

spacings (0.015 cm < L < 0.165 cm). Note that logarithmic scales are employed both for frequency (from 0.3 to 50 cm
1) and intensity, allowing a direct check of the scaling properties. The peak at about 30 cm
1 is attributed to the acoustic vibrations of the basic units, a spheroidal cluster of porous silica with a size (diameter), a, of 1.2–1.6 nm (see “Raman Spectroscopy of Nanocrystals”). The higher limit, ξco, and the related ωco1 depend on the density of the aerogel. At frequencies ω < ωco1, phonon-like acoustic vibrations are present and a steeper slope is observed in the frequency dependence of the Raman intensity. For the lightest aerogel, the linear behavior extends over 1.5 decades of frequency. Its low-frequency extension is limited by the measured frequency range, in agreement with the value ωco1  0.02 cm
1, derived from Brillouin data (Courtens et al. 1987, 1988). For the heaviest aerogel, Brillouin data gave ωco1  0.3 cm
1, a value which roughly corresponds to the onset of the fractal behavior. Similar results were obtained for base-catalyzed

38

Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies

1163

aerogels, the main difference being the larger size of the building blocks, a ffi 3.2 nm, and a particle peak at about 15 cm
1. In Fig. 3, the susceptibility I(ω)/(n + 1) is reported. Therefore, from Eq. 7 and from the hypothesis that C(ω)  ωx, the observed slope y =
0.36 should be given by y ¼ x þ d~
2: The independent measurement of the fractal dimension from SANS experiments, D = 2.4 (Vacher et al. 1988), and of the spectral dimension by INS experiments, d~ ¼ 1:3  0:1 (Vacher et al. 1989, 1990; Courtens et al. 1990), allowed the determination of the scaling exponent of C(ω): x = 0.34. Different models have been proposed for the calculation of C(ω), based on the scaling properties of the local strain in fractals (Boukenter et al. 1986; Tsujimi et al. 1988; Alexander 1989; Alexander et al. 1993; Mazzacurati et al. 1992; Duval et al. 1993). Molecular dynamics simulations on model systems, n-dimensional percolators, were also performed by using different Raman scattering mechanisms, DID or bond polarizability (BP) (Montagna et al. 1990; Stoll et al. 1992; Mazzacurati et al. 1992). Also, due to the size limitations of these simulations, a definitive assessment of the scaling properties of C(ω) was not reached. The situation is particularly complicated for DID scattering mechanisms, since the vibrations and the associated strains can propagate only along the disordered connections of the fractal, but electric fields can propagate also in the free space of the porous structure (Alexander 1989; Montagna et al. 1990; Mazzacurati et al. 1992). The microscopic structure of the system and the actual scattering mechanism, DID or BP, seem to determine the shape of the Raman spectrum at least at the same extent as do the macroscopic properties described by the fractal parameters (Mazzacurati et al. 1992). Raman and Brillouin studies were extended to silica aerogels produced with different procedures, having different microstructures, connectivities, and thus different vibrational dynamics. In particular, a series of samples were prepared with different densities ρ and thus different crossover lengths ξco, but with nearly identical structures on fractal scales, i.e., for distances smaller than ξco. SANS measurements showed a fractal behavior, allowing one to obtain the fractal dimension D and the fractal range between a and ξco (Vacher et al. 1988). Observations in direct space, by electron microscopy, were in good agreement with SANS results (Rousset et al. 1990; Courtens and Vacher 1992). Furthermore, INS experiments and neutron spin echo spectroscopy allowed the measurement of the density of vibrational states and thus d~ (Schaefer et al. 1990; Conrad et al. 1990; Courtens et al. 1990; Vacher et al. 1990). A full agreement was found between the values of D, d~ , and ωco1, derived from Brillouin and Raman scattering experiments, and those obtained by SANS, INS, and electron microscopy experiments. The structure strongly depends on the preparation conditions. Silica gels obtained with basic catalysis produce fractal structures composed of connected spheroidal silica units of nanometric size (a = 1–2 nm). In gels obtained with acid catalysis, the structural units are the SiO4 tetrahedra. Gels reacted without the addition of a catalyst produce aerogels with a higher fractal dimension (D  2.4) than gels obtained with basic catalysis (D  1.8) (Courtens and Vacher 1992). The first measurement of d~ by Brillouin and INS experiments in a neutrally reacted series of aerogels with D = 2.46, d~ = 1.3, was very

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M. Montagna

close to the universal value d~ = 4/3 suggested for percolating networks (Alexander and Orbach 1982). However, further measurements in base-catalyzed series provided d~ ¼ 1:4, D = 1.89, showing that the fracton dimension d~ is not a universal quantity, but is an additional dimension, which describes the connectivity within the fractal system (Courtens and Vacher 1992). Furthermore, the density of vibrational states obtained by INS experiments showed that d~ is not really constant for frequencies between ωco1 and ωco2, but is higher at high frequencies (Vacher et al. 1989, 1990; Courtens et al. 1990). These two different regimes were associated to bending in the low frequency range and to stretching in the high frequency range. This observation confirmed the prediction of Feng that the elasticity of tenuous materials is dominated by stretching at short scales and by bending at longer scales (Feng 1985). The scaling of the depolarized Raman spectra with a single power y in the two frequency ranges, where bending or stretching dominate, can probably be explained by a stronger Raman activity of bending over stretching vibrations (Courtens and Vacher 1992). Low-temperature specific heat and thermal conductivity measurements, between 0.05 and 20 K, confirmed that in silica aerogels three different temperature regimes are present. These three regimes correspond to phonon, fracton, and particle mode vibrations, in increasing order of temperature (Bernasconi et al. 1992). These studies were subsequently extended to partially densified aerogels, prepared by heat treatment and hydrostatic or uniaxial pressure. Changes in the spectra were correlated with densification-induced changes in structure and connectivity of the aerogels. By combining SAXS and SANS, INS, Raman, and Brillouin results, it ~ ωco1 and ωco2, ξco, and a and was possible to obtain the values of the parameters D, d, to study their evolution with densification. Heat treatment induces matter displacements at all scales including the particle scale, a, whereas the effect of pressure essentially affects the long-range structure of fractal aggregates. In base-catalyzed aerogels, the mean size, a, of the spheroidal silica particles increases with thermal treatment. This is shown by the shift of the Raman peak, the peak in the highfrequency side of Fig. 3, to low frequencies, and is confirmed by SAXS measurements (Anglaret et al. 1998). Progressive densification by thermal treatment reduces ξco (increases ωco1), while maintaining constant D and d~ . In these base-catalyzed silica aerogels, D  1.75 and d~  1:1, values quite lower than those of neutrally reacted aerogels (Anglaret et al. 1995). Brillouin measurements show that the elastic modulus increases strongly during sintering, while the attenuation decreases, which is coherent with a larger connectivity in the solid network. Viscous flow sintering creates new siloxane bonds, eliminates pores, and, as expected, stiffens the aerogel (Calas et al. 1998; Caponi et al. 2003). Partially densified aerogels, prepared by hydrostatic or uniaxial pressure, have a different behavior. Brillouin measurements show that the elastic modulus decreases and the attenuation increases in the low-pressure regime (Calas et al. 1998). This is attributed to a plastic deformation with breakage of links between clusters during compression. For higher pressure, the density increase is accompanied by stiffening, suggesting that condensation occurs more than link breakage. Raman and Brillouin data show that d~ increases with

38

Characterization of Sol-Gel Materials by Raman and Brillouin Spectroscopies

1165

pressure-induced densification. This is attributed to a strengthening of the bonds at the particle scale, which produces an increase of the density of vibrational states at high frequencies (Anglaret et al. 1998). The application of uniaxial pressure produces elastically isotropic samples, very similar to those obtained by hydrostatic pressure (Levelut et al. 1998). The macroscopic effect of compaction is highly anisotropic. However, the new siloxane bonds created by the asymmetric compaction are not oriented. They are created in all directions by structural rearrangements at the microscopic scale. In aerogels, the measured attenuation, given by the Brillouin line width, is mainly caused by the structural disorder present in the amorphous structure (Courtens et al. 1987). This process is almost temperature independent. On the contrary, in vitreous silica, the attenuation is mainly due to the presence of thermally activated processes, such as relaxations and two-level systems (Vacher et al. 1997). These mechanisms have been proposed to explain the strong increase, about one order of magnitude, of the Brillouin line width in v-SiO2, from liquid helium to room temperature, where Γ  75 MHz  0.0025 cm
1 (Vacher and Pelous 1976). This occurs for Brillouin scattering of visible light, where the phonon frequency is of the order of 1 cm
1 and the exchanged q vector is of the order of 0.04 nm
1. In inelastic X-ray scattering experiments in silica, where ωph  30 cm
1 and q  1 nm
1, the Brillouin line width is quite large, with Γ comparable to the peak frequency (Foret et al. 1996; Benassi et al. 1996; Pilla et al. 2000, 2002; Ruocco and Sette 2001). It is temperature independent and thus attributed to the structural disorder, which strongly attenuates the phonon propagation at these frequencies. It increases with the q vector as Γ / qα, with α  2 (Ruocco and Sette 2001). Therefore, in silica, phonon attenuation is dominated by the disordered structure (static attenuation mechanisms) at high frequency and by glass relaxation (dynamic mechanism) at ωph  1 cm
1. The frequency region in between, where the two mechanisms should give comparable effects, is not covered by experimental Brillouin facilities. However, phonon attenuation produced by structural disorder is higher in porous systems, as partially densified aerogels and xerogels. In these systems, the two mechanisms of phonon attenuation are comparable in visible and UV Brillouin scattering (Caponi et al. 2003, 2004). It has been found that for pore sizes smaller than about 8 nm, the acoustical attenuation is the same as in v-SiO2 at room temperature. The attenuation was therefore attributed to dynamical processes. For larger pore sizes, the Brillouin line width is larger and it becomes nearly temperature independent and strongly increases with the mean pore size. The main mechanism of attenuation is now phonon scattering by the disordered porous structure. Dealing with these porous hygroscopic systems, one needs to take care of removing water by high-temperature thermal treatment in order to isolate the dynamics of dry systems. In samples left in air for days, water adsorbed in the porous structure produces a decrease of the sound velocity and an increase of attenuation (Terki et al. 1999; Fontana et al. 1999). Furthermore, water dynamics gives strong low-frequency Raman scattering (Terki et al. 1999; Cicognani et al. 1999).

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Densification of Silica Xerogels Figure 4 shows the VV-polarized Raman spectra of three xerogels, heated at different temperatures, together with that of v-SiO2. The three samples were prepared using as starting solution a mixture of tetramethyl orthosilicate (TMOS), methanol, deionized water, and nitric acid, in the molar ratio 0.06:0.35:0.55:0.04, respectively. The solvent was removed by evaporation in air, at room temperature, for several weeks. The progressive densification of silica gels was obtained by thermal treatment in air, for 72 h, at several temperatures (Tt), reached with a rate of 0.1 C/min (Caponi et al. 2002). The spectra of Fig. 4 cover the frequency range of the O–H stretching vibrations. The porous samples heated at 600 C, with a density of 1300 kg/cm3, show a broad band, extending from about 3200 to about 3800 cm
1. This band is due to the O–H stretching vibration of the “physical” and “chemical” water in the pores of the gel. Physical water, i.e., the residual water in the pores of the structure, produces a broad band, which is typical of a liquid system. Chemical water, i.e., a residual layer of water on the pore surface, produces the weak and sharper structures superimposed on the broad band. Different vibrational frequencies Fig. 4 VV-polarized Raman spectra of xerogels annealed at different temperatures Tt in the high-frequency region of the O–H stretching vibrations. For comparison, the spectrum of vitreous silica is also reported (Caponi et al. 2002)

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correspond to different configurations of bonds of the chemical water with the Si–OH units and with the neighbor water molecules at the pore surface. The sharp peak, centered at about 3740 cm
1, is attributed to the O–H stretching vibrations of free silanol groups on the pore surface (Bertoluzza et al. 1982; Gottardi et al. 1984). It is much sharper than the bands of physical and chemical water, because the Si–OH groups at the pore surface are only weakly perturbed by site-sensitive interactions. After annealing at 860 C (ρ = 1800 kg/m3), the broad band disappears. The sharp peak of the free silanol groups dominates the spectrum. Residual weak bands, due to chemical water, are also present. The thermal treatment produced a nearly complete evaporation of water molecules, but the temperature was not sufficiently high to produce densification by viscous sintering (Brinker and Scherer 1990). Therefore, the structure maintained a porous structure. The sample annealed at 900 C is a densified glass with negligible porosity. The peak of free silanol groups is no longer present, and a broader band at about 3650 cm
1, also present in the spectrum of v-SiO2, appears. It is attributed to hydrogen-bonded silanol groups (Gottardi et al. 1984). The O–H groups, now embedded in the silica matrix, have sitedependent vibrations, and a Raman peak broader than that of the free silanol groups is observed. Figure 5 shows the VV spectra of some xerogels heated at different temperatures, together with that of v-SiO2, in the whole frequency range of vibration of the silica network. We can see that the major features present in v-SiO2 are also observed in the xerogels. The intense broad band at about 400 cm
1 is attributed to stretching–bending vibrations of the silica network (Galeener 1985). Its width is related to the disorder-induced distribution of angles in the Si–O–Si units, which connect the SiO4 tetrahedra. The band is sharper in the xerogels than in v-SiO2, since the open porous structure has a lower number of constraints than the harder compact silica network. The two sharp peaks of silica at about 490 and 605 cm
1, known as D1 and D2 defect bands, are attributed to local defects with four- and threefold rings, instead than the normal sixfold rings of the silica network (Galeener 1985). In the spectra of xerogels, these bands appear more intense than in silica. Their line shapes are different from those of silica, since the two peaks are broader and slightly shifted in frequency, the D1 band being centered at 478 cm
1 (Walrafen et al. 1985, 1986; Mulder and Damen 1987; Brinker et al. 1990). The Raman spectra clearly show that three- and fourfold rings are favored on the pore surface. Note that the densified xerogel, annealed at 900 C, shows D1 and D2 bands less intense than v-SiO2, indicating a less defective structure. The band at about 800 cm
1 and the two weak bands at 1050 and 1200 cm
1 are attributed to symmetric and antisymmetric Si–O stretching vibrations, respectively. These bands have practically the same shape in the xerogels and in v-SiO2, indicating that the building blocks, i.e., the SiO4 tetrahedral structures, are the same. The relatively sharp band at 970 cm
1, whose intensity decreases with thermal treatment, is due to Si–OH stretching vibrations of silanol groups (Bertoluzza et al. 1982; Gottardi et al. 1984). The peak at about 50 cm
1, present in the spectra of v-SiO2 and of the densified xerogel, is the boson peak, a common feature of the room temperature Raman spectra of all glasses. Its shape appears more clearly in Fig. 6, which shows the HV Raman spectra in the

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Fig. 5 VV-polarized Raman spectra of xerogels annealed at different temperatures Tt. For comparison, the spectrum of vitreous silica is also reported (Caponi et al. 2002)

frequency region below 100 cm
1. HV Raman spectra are more useful for the study of the boson peak, which is depolarized with comparable intensities in HV and VV polarizations. In fact, the intense band of Fig. 5, centered at 400 cm
1, but having a long tail toward low frequencies, is almost completely polarized and thus does not appear in HV polarization. Figure 6 shows that the structure of these xerogels abruptly changes in the 860 C < Tt < 875 C temperature range, with the sudden appearance of the boson peak. This result is attributed to the peculiarity of the densification process. The viscous sintering, with the elimination of pores and dangling Si–OH bonds, occurs at a well-defined temperature, producing a glass with a structure close to that of v-SiO2. In the non-densified xerogels, the Raman spectra show a relatively sharp band centered at about 16 cm
1, for the xerogel with Tt = 800 C, and at about 18 cm
1, for the xerogel with Tt = 860 C. In the samples treated below 800 C, it is not possible to observe the low-frequency bump. The strong quasi-elastic scattering is due to vibrational dynamics of the porous fractal-like system, but probably also to residual water inside the pores. The low-frequency bumps are attributed to surface

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Fig. 6 HV-polarized Raman spectra in the low-frequency region of xerogels annealed at different temperatures, Tt, and of vitreous silica: (a) several heat treatment temperatures and vitreous silica, with spectra normalized to the intensity of the band at 800 cm
1; (b) spectra normalized to the intensity of the low-frequency band (Caponi et al. 2002)

vibrations of the particle–pore structure. The peak frequency increases with the annealing temperature, as the mean pore size, measured by SAXS experiments, decreases (see “Raman Spectroscopy of Nanocrystals”). The Raman data in the low frequency range of Fig. 6 show that the vibrational dynamics are almost the same for the densified xerogels and v-SiO2; the boson peak is centered at the same frequency and has a similar spectral shape. An excess of scattering in the low-frequency part of the boson peak is present in the densified xerogel samples heated at 875 C and 900 C, indicating the presence of a small amount of residual porosity, to which low-frequency Raman scattering is very sensitive. The density and SAXS measurements on the sample treated at 900 C show a small residual porosity with large pore size (6 nm) and indeed small internal surface area. Low-temperature specific heat and Brillouin scattering measurements confirm the Raman results. Important specific heat excess is observed for temperatures lower than about 5 Kin the non-densified xerogels, whereas the specific heat of the sample annealed at 900 C is identical to that of v-SiO2 (Caponi et al. 2002). The sound velocity of the xerogel increases with the annealing temperature, suddenly reaching that of v-SiO2 for Tt 875 C (Bartolotta et al. 2001). The appearance of the boson peak in the Raman spectra is now considered a signature of the occurred densification (Mariotto et al. 1988a; Armellini et al. 1998). This is

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particularly important when densification is needed, but at the lowest possible annealing temperature, in order to avoid or reduce crystallization effects (Bouajaj et al. 1997; Montagna et al. 2003; Zampedri et al. 2003). It has been observed that silica xerogels activated with rare earth ions have densification temperatures that depend on the content and nature of the doping ions. Doping with Tb3+ increases the transition temperature, whereas the opposite occurs by doping with Pr3+ (Pucker et al. 1998; Armellini et al. 1998).

Characterization by Waveguided Brillouin Spectroscopy As discussed in “Inelastic Light Scattering,” Brillouin scattering in bulk glasses produces sharp peaks, whose line width is determined by the dynamical mechanism of phonon attenuation. In fact, the exchanged q is precisely determined by the geometry of the experiment, since the scattered wave is collected from a large volume, with a typical size of 50 50 5000 μm3, uniformly illuminated by the plane wave of the laser beam. On the contrary, an exciting light beam, confined in a planar waveguide, does not have a single well-defined wavevector. In a thick waveguide, a ray-optic approach can be used for the description of wave propagation, as shown in Fig. 7. The light propagates in the z direction and the plane waves have a zigzag path in the x–z plane, undergoing total internal reflection at the boundary interfaces of the waveguide. The laser beam, polarized along the y direction, can be injected into the waveguide by prism coupling and propagates only at discrete values of the angle ϑ in one of the transverse electric (TEm) modes. The m value (m = 0, 1, 2,. . ., mmax) gives the number of nodes of the electric field, a stationary wave, along the x direction. As shown in Fig. 7, when the scattered light is collected from the front surface of the waveguides, two exchanged q vectors are present. A simple model, which takes into account the relative phases of the two scattered waves, shows that two main peaks due to longitudinal phonons are present, Fig. 7 Wave propagation in the planar waveguide. q1 and q2 are the exchanged wavevectors of the scattered light in the zigzag paths. In TE modes, the electric field is along the y direction, the light propagates along the z direction, and x is perpendicular to the plane of the waveguide in the direction of the scattered wave. ns and ng stand for the refractive indices of the substrate and waveguide, respectively (Chiasera et al. 2003a)

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apart from the case of the TE0 excitation, where a single peak is observed (Montagna et al. 1998). The energy separation between the two peaks increases with the mode index. This model neglects the contribution to the scattering coming from the evanescent field in the substrate and considers the waveguide as a homogeneous film with constant refractive index. A single fit parameter, i.e., the longitudinal sound velocity, is used to calculate the m + 1 spectra obtained by exciting the different modes of the waveguide. In graded-index waveguides, m – 1 peaks of comparable intensities are observed. Figure 8 shows the Brillouin spectra of a silica–titania graded-index waveguide. The sample was deposited by a dip-coating technique (Zampedri et al. 2003). The graded-index SiO2–TiO2 planar waveguide was obtained by subsequently depositing 35 layers of TiO2(0.08)–SiO2(0.92) composition, 25 layers of TiO2(0.16)–SiO2(0.84), and finally 23 layers of TiO2(0.24)–SiO2(0.76) composition (molar fractions in parenthesis), on a silica substrate. After each dip, the films were annealed in air at 900 C for 30 s. After every ten dipping cycles, the films were heat treated at 900 C for 2 min. Finally, the waveguides were submitted to a further heat treatment at 1000 C, with a heating rate of 20 C/min from 600 C to 1000 C. The waveguide was characterized by m-line spectroscopy. It supports five TE and TM modes at 543.5 nm and four at 632.8 nm. The refractive index profile, calculated at 514.5 nm by extrapolation of the data at 543.5 and at 632.8 nm, is shown on the upper left side of Fig. 8. The calculated profiles of the squared electric field show that the different modes selectively excite the different layers of the waveguide; the TE0 mode is practically completely confined in the third (from the substrate) layer with TiO2(0.24)–SiO2(0.76) composition, while the TE4 mainly occupies the first layer, with TiO2(0.08)–SiO2(0.92) composition. The Brillouin spectra in the longitudinal phonon spectral region are shown on the right-hand side of Fig. 8, together with the results of a numerical model, which considers the spatial distribution of the exciting field in the mode, a simple spatial dependence of the elasto-optic coefficients, through the value of the refractive index, and neglects the refraction of phonons (Chiasera et al. 2003a). A single fit parameter, i.e., the sound velocity, is necessary to obtain the calculated spectra. For the TE0, TE1, and TE2 spectra, which show one, two, and three peaks, respectively, the agreement is good, even if the observed intensity of the higherfrequency components in TE1 and TE2 spectra is lower than expected and the observed line width is slightly larger than expected. In the TE3 spectrum, only three peaks are observed, instead of the four calculated ones. TE4 shows a single strong peak with shoulders, partially reproduced by the calculations. In any case, a general agreement is present in all spectra, sufficient to determine with high accuracy the longitudinal sound velocity. Within the experimental error, we used the same value υL = 5.9 km/s for the four excitations with m  3. These four modes involve, with different weights, the two external layers with the compositions TiO2(0.24)–SiO2(0.76) and TiO2(0.16)–SiO2(0.84). We should conclude that the longitudinal sound velocity is nearly the same for the two compositions. A quite weak dependence of the sound velocity on the titania content was already observed (Montagna et al. 1998). On the contrary, the TE4 spectrum, with most excitation in the internal layer with TiO2(0.08)–SiO2(0.92) composition, shows a lower value of

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Fig. 8 Upper left-hand side frame: refractive index profile at 514.5 nm of a three-layered SiO2–TiO2 planar waveguide. Left-hand side column: calculated squared electric-field patterns of the five TEm modes. Right-hand side column: Brillouin experimental spectra (open circles), calculated spectra (dotted line), and convolution of the calculated spectra with the instrumental response (solid line). The longitudinal sound velocity used in the fit is υL = 5.9 km/s, for m = 0, 1, 2, and 3, and υL = 5.75 km/s, for m = 4 (Chiasera et al. 2003b)

the longitudinal sound velocity, υL = 5.75 km/s. This is attributed to a residual porosity of this internal layer, which did not undergo full densification (Chiasera et al. 2003b). The result was confirmed by waveguided Raman spectroscopy. The Raman spectrum obtained by exciting the waveguide in the TE4 mode shows residual porosity. In particular, a weak band due the stretching vibrations of free Si–OH groups appears. On the contrary, the Raman spectra obtained by exciting in the TE0, TE1, and TE2 modes show that the two external layers are completely

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densified. This was probably due to densification of the external layers occurring at a lower temperature than that of the internal first one, which was indeed not allowed to freely expel its reaction products.

Raman Spectroscopy of Glass Ceramics Raman spectroscopy is a very powerful technique to follow the devitrification process induced by thermal treatments, which produce glass ceramics. A good example is that of silica–titania xerogels (Almeida and Christensen 1997; 1998; Strohhöfer et al. 1988; Bersani et al. 1998a, b; Bersani et al. 1988; Karthikeyan and Almeida 2000; Montagna et al. 2003; Zampedri et al. 2003). Figure 9 shows the Raman spectra of a series of TiO2(x)–SiO2(1 – x) waveguides obtained by dip coating and annealing at 900 C. The waveguides are activated with 1 mol% Er3+, they have a thickness between 1.4 and 2.2 μm, and they support a TE and a TM mode at 1.5 μm, plus two or three modes in the visible. The waveguiding configuration was used for the measurements. The Raman spectra were collected in VV polarization, by exciting by prism coupling the TE0 mode with an Ar+-ion laser operating at 457.9 nm and detecting the scattered light from the front of the waveguide (Zampedri et al. 2003). All spectra show the characteristic peaks of the silica. The bands at about 950 and 1100 cm
1 are assigned to the vibrations of mixed Si–O–Ti linkages (Best and Condrate 1985). All spectra show an initial crystallization process, evidenced by the Fig. 9 VV-polarized Raman spectra of TiO2(x)–SiO2(1
x): Er3+ waveguides with different TiO2 contents. The VV spectrum of SiO2 glass is also shown (Zampedri et al. 2003)

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Fig. 10 VV-polarized Raman spectra of SiO2(0.8)–TiO2(0.2) waveguides. The labels indicate the temperature ( C) and time (min) of annealing. Excitation was at 514.5 nm, by prism coupling, in the TE0 mode (Montagna et al. 2003)

structure in the region between 150 and 350 cm
1, attributed to optical vibrations of TiO2 crystals (Moret et al. 2000), and by the low frequency peak, due to the acoustic vibrations of titania nanocrystals (Montagna et al. 2003), which partially overlaps with the boson peak of the glass at about 40 cm
1. They show that devitrification of the silica–titania film occurs and this effect is more important for the waveguides with 15, 20, and 24 mol% of TiO2. Pure glassy films are obtained by annealing at 700 C after each dip. In any case, the successive annealing processes at higher temperature, necessary for a full densification of the xerogel, produced some degree of devitrification (Zampedri et al. 2003). Figure 10 shows the VV Raman spectra of thinner SiO2(0.8)–TiO2(0.2) waveguides, with a thickness of about 0.4 μm, treated at different temperatures (Montagna et al. 2003). The spectrum of the waveguide heated at 700 C shows the typical bands of silica–titania amorphous network, with no visible contributions from a crystalline phase. After annealing at 800 C, sharp peaks appear, superimposed on the broadband spectrum of the glass. These peaks become more and more intense with increasing annealing temperature and time. They are attributed to optical vibrations of crystalline TiO2. For Ta 1200 C, only the characteristic peaks of the anatase phase are observed (Haro-Poniatowski et al. 1994; Moret et al. 2000), with a sharp and intense peak at about 141 cm
1, attributed to Eg vibration and three other peaks at about 394 cm
1 (B1g), 513 cm
1 (A1g + B1g), and 635 cm
1 (Eg). The frequencies of the peaks are smaller than those

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of the bulk crystal (Haro-Poniatowski et al. 1994), because of phonon confinement in the nanoparticles (see “Raman Spectroscopy of Nanocrystals”). For Ta in the range 800–1200 C, the Raman spectra indicate that the nanocrystals are a mixture of anatase and brookite phases (Moret et al. 2000). The anatase content seems to increase progressively with the annealing temperature. At very low frequency, a structured band appears in Fig. 10. It increases in intensity and progressively shifts toward low frequency with the annealing temperature. It is assigned to the acoustic vibrations of the titania nanocrystals, and its frequency position allows the determination of the particle size, as we will discuss in detail in the next section.

Raman Spectroscopy of Nanocrystals Conventional methods for measuring nanoparticle size are transmission electron microscopy, in which the particles are directly imaged, and X-ray diffraction, in which the particle size is inferred from the width of the diffraction lines, using the Sherrer method. Raman scattering of optical and, in particular, of acoustic vibrations is a simple, fast technique for obtaining the size distribution of nanoparticles. Phonon confinement in the nanoparticles produces shift and broadening of the Raman lines of optical phonons (Richter et al. 1981; Tiong et al. 1984; Campbell and Fauchet 1986). The momentum conservation rule of Eq. 5 does not hold for small crystalline particles with qa < 1, where a is the particle size. A phenomenological model, i.e., the spatial correlation model, can account for the frequency shift. Confined optical phonons with any kph contribute to the Raman scattering (Campbell and Fauchet 1986):    C q, kph 2 I ð ωÞ /  dkph ; 2 ω
ωph þ ðΓ0 =2Þ2 ð

(10)

where Γ0 is the width of the Raman line in the bulk crystal. Since the q vector of visible light is much smaller than the kph of phonons in the nanocrystal (q  0), a Gaussian function is usually taken for the Fourier coefficient of the confinement function: !    k 2 a2 C q, kph 2 ¼ exp
ph : 16π 2

(11)

The phonon dispersion relation ω(kph) of the bulk modes is usually taken. For most crystalline solids, the phonon dispersion relation has a maximum at the center of the Brillouin zone, and one observes a low-frequency shift of the Raman line with decreasing crystalline size, but the opposite behavior can also occur (Nemanich et al. 1981).

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The technique has been applied to the size determination of TiO2 and PbTiO3 nanoparticles in films deposited by sol-gel (Bersani et al. 1998a, b; Bersani et al. 1988). The mechanisms that give rise to the broadening and the shifts of the Raman peaks of titanium dioxide and lead titanate nanocrystals prepared by sol-gel are different. Phonon confinement and oxygen deficiency are competitive mechanisms in TiO2 obtained by different sol-gel preparations, whereas pressure effects on the nanocrystals predominate in ferroelectric PbTiO3. For CdS nanocrystals in silica xerogels, no shift of the Raman line was observed (Capoen et al. 2001). It was suggested that the theoretically predicted redshift due to the phonon confinement may be hindered by a blueshift. This second effect would be caused by strain acting on the surface of the nanocrystals (Shiang et al. 1993). In fact, phonon confinement is one possible cause of shift and broadening of the Raman line, but there are other causes, as the presence of a size distribution. Stress also causes a shift of the line. For these reasons, a reliable measurement of the size of the nanocrystals is rarely obtained from the line shape of optical phonons. On the contrary, a much more precise measurement of the nanocrystal size can be obtained from Raman scattering of the acoustic vibrations. After the first works on spinel nanocrystals in cordierite glasses (Boukenter et al. 1986) and on silver colloids in alkali halides (Mariotto et al. 1988b), low-frequency Raman scattering from symmetric and quadrupolar acoustic vibrations of nanoparticles has become a nondestructive method to determine the size of the particles. A peak in the range 5–50 cm
1 was observed in many composite systems containing metallic, insulator, or semiconductor nanoparticles (Fujii et al. 1991, 1996; Ferrari et al. 1996, 1999; Ceccato et al. 2001; Tikhomirov et al. 2002; Montagna et al. 2003; Ivanda et al. 2003). The size of the nanoparticles is derived from the energy of the peak, since the frequency of all modes scales as the inverse of the linear dimension of the particles. The acoustic vibrations of an elastic homogeneous sphere with a free surface are classified as spheroidal and torsional modes (Lamb 1882). Torsional vibrations involve only shear deformations and are not Raman active (Duval 1992). Spheroidal modes involve both shear and stretching motions and produce radial displacements. They are characterized, following the symmetry of the sphere, by three labels (l, m, p). The symmetric l = 0 (m = 0) spheroidal modes are purely radial with spherical symmetry. At higher l values, angular corrugations appear. l measures the number of wavelengths along a circle on the surface. A third index, p = 1, 2, . . . labels the sequence of modes in increasing order of frequency and radial wavevector at fixed angular shape (l, m). The quantity p
1 measures the number of nodes of the vibrations in the radial direction. The fundamental p = 1 mode is called surface mode, its overtones ( p > 1) being called inner modes. Only the symmetric (l = 0, m = 0) and the quadrupolar (l = 2,
2 < m < 2), with fivefold degeneracy, are Raman active (Duval 1992). The frequencies of all modes scale as the inverse radius of the sphere, allowing the introduction of the dimensional quantities: hR ¼ ω0 R=υL ,

kR ¼ ω2 R=υT ;

(12)

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Fig. 11 Dimensional frequencies of the l = 0 (hR) and l = 2 (kR) surface modes of a free sphere, as a function of the ratio between the longitudinal and transverse sound velocities, υL/υT

for the symmetric and for the quadrupolar modes, where υL and υT are the longitudinal and transverse sound velocities, respectively. hR and kR depend only on the ratio of the longitudinal and transverse sound velocities υL/υT. This dependence is shown in Fig. 11, for the surface modes (p = 1). If the sound velocities of the particles are known, one can obtain a first rough value of the particle size from the position of the l = 0 and/or l = 2 peak, by using the relations (12). The symmetric l = 0 modes give polarized Raman spectra (IVH = 0); the quadrupolar l = 2 modes give depolarized spectra. Therefore, on the basis of the depolarization ratio, IVH/IVV, the Raman peaks can be assigned to symmetrical or quadrupolar vibrations. The depolarization ratio IVH/IVV of the quadrupolar modes and the relative efficiencies of the quadrupolar and symmetric modes are system dependent, since they depend on the microscopic structure and scattering mechanism (Montagna and Dusi 1995). In silver nanoparticles, only the depolarized quadrupolar l = 2 vibrations are Raman active (Fujii et al. 1991; Ferrari et al. 1995, 1996, 1999; Palpant et al. 1999). This occurs because the symmetric l = 0 vibrations are not Raman active in crystals with a cubic Bravais lattice (Montagna and Dusi 1995). On the contrary, for CdS nanocrystals, the symmetric l = 0 modes dominate the Raman spectra, the contribution of the quadrupolar vibrations being relatively weak (Saviot et al. 1998; Ivanda et al. 2003). In other systems, as Ga2O3 and TiO2, quadrupolar and symmetric vibrations have comparable intensities (Ceccato et al. 2001; Montagna et al. 2003). Figure 12 shows the spectra of the acoustic vibrations of TiO2 nanoparticles, grown by thermal treatment of silica–titania waveguides obtained by dip coating (Montagna et al. 2003). The l = 2 surface vibrations, which are active both in VV and VH polarizations, produce the lowest frequency peak. The intense peak at higher frequency, present only in VV spectra, is attributed to the l = 0 surface mode. The weaker peak at higher frequency, at about 35 cm
1 in sample s1200(60) (annealed at 1200 C for 60 min) and at about 23 cm
1 in sample s1300(30), is due to an inner l = 0 mode (p = 2), a shorter wavelength symmetric mode with a node in the radial wave function. In

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Fig. 12 Low-frequency VV and HV Raman spectra of SiO2 (0.8)–TiO2 (0.2) waveguides. The labels indicate the temperature ( C) and time (min) of annealing. Excitation was at 514.5 nm, by prism coupling in the TE0 mode (Montagna et al. 2003)

general, the intensities of the peaks of a sequence ( p = 1, 2, . . .) rapidly decrease with p, especially in the l = 2 case (Montagna and Dusi 1995), so that only the surface p = 1 mode can be easily detected. All peaks shift toward lower frequencies as the annealing temperature increases, showing the progressive increase of the mean size of the nanocrystals. The Raman spectra were fitted by considering that the line width of the peaks has two main sources: the homogeneous broadening, due to the interaction of the vibrating particle with the surrounding glass (Montagna and Dusi 1995), and the inhomogeneous broadening, due to a distribution of the particle sizes. For annealing temperatures between 900 C and 1300 C, the mean size increases from about 4 to 20 nm in diameter. Crystallites are present even after annealing at 800 C, but their size cannot be well evaluated from the low-frequency Raman spectra, because the relative scattering is weak and not well resolvable from the boson peak of the glass.

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The mean particle size obtained by the Raman spectra compare well with those obtained from the line width in X-ray diffraction measurements, for particle sizes larger than about 8 nm. For smaller particles, Raman data give sizes larger than those obtained by X-ray measurements (Montagna et al. 2003). Low-frequency Raman scattering from acoustic vibrations was also employed to determine the size of CdS nanocrystals in silica xerogels (Othmani et al. 1992).

Conclusions The survey presented in this chapter shows that Raman and Brillouin spectroscopies in sol-gel-derived materials cover a quite vast domain of investigation, from the basic glass science to the characterization of materials produced for many different applications. The evolution of Raman instrumentation is now allowing extending Raman spectroscopy from the laboratory to industry, as monitoring process for quality control.

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Small-Angle X-ray Scattering by Nanostructured Materials

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Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Basic Theory . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . General Equations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Small-Angle Scattering by Nanoscopic Two-Phase Systems: Porod’s Law . . . . . . . . . . . . . . . Small-Angle Scattering by Spatially Uncorrelated Nanoparticles: Guinier’s Law . . . . . . . . Dilute Sets of Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Spherical Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Application (Example 1): PbTe Nanocrystals Embedded in a Silicate Glass . . . . . . . . . . . . . Application (Example 2): Clustering of Colloidal ZnO Nanoparticles . . . . . . . . . . . . . . . . . . . . Concentrated Sets of Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Spatially Correlated Spherical Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Application: Fe-Doped Organic–Inorganic Hybrid Nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . Fractal Structures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Small-Angle Scattering by Fractal Structures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Applications: Aggregation in Zirconia-Based Sols and Gels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nanophase Separation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . General Considerations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Phase Separation and Dynamical Scaling Property . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Application (Example 1): Sintering of SnO2-Based Xerogels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Application (Example 2): Dynamical Scaling of Zirconia-Based Fractal Structures . . . . . Grazing Incidence Small-Angle X-Ray Scattering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Basic Concepts . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Example of Application: Nanostructure of Thin Films Supported by Si Wafers . . . . . . . . . . Final Remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Appendix: Experimental Issues . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Basic Comments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Choice of Sample Thickness . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Subtraction of Parasitic Scattering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Correction of Smearing Effects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Determinations of SAXS Intensity in Relative and Absolute Units . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

This chapter contains the basic theory of small-angle X-ray scattering (SAXS) and its applications to low-resolution studies of nanostructured materials. The primary purpose is to explain how to obtain structural information from simple systems whose low-resolution structure can be described by a two-electron density model, consisting of either homogeneous nanoparticles embedded in a (solid or liquid) medium with constant electron density or two-phase bicontinuous systems. The presented SAXS theory and the examples of applications refer to different procedures for determinations of geometrical parameters associated to nanoparticles or clusters in dilute solution, spatially correlated nanoparticles, and more general two-phase systems, namely, particle radius of gyration, interface area, size distribution, fractal dimension, and interparticle average distance. Other described applications are in situ SAXS studies of mechanisms involved in transformation processes leading to nanostructured materials such as those occurring in nanophase separation and along the successive steps of sol-gel routes. One section is dedicated to present the basic concepts and describes an application of grazing incidence small-angle scattering (GISAXS), which allows for studying nanostructured thin films and thin layers located close to the external surface of solid substrates. Most of the reported applications refer to nanostructured materials obtained by sol-gel processing and are based on experimental results published by the author and collaborators.

Introduction This chapter describes the basic theory of small-angle X-ray scattering (SAXS) and reports a number of examples of application of this experimental technique to low-resolution structural investigations. Several examples also show how SAXS is applied to the characterization of transformation mechanisms in different nanomaterials. Sol-gel processing starts from colloidal particles in liquid solution and often leads to solid materials with interesting properties. Along all steps of sol-gel transformations the nanoscopic nature of the structure is preserved. This chapter includes several applications of SAXS technique to in situ characterizations of precursor systems starting from liquid sols up to final solid nanostructured materials. The basic process of the scattering of X-rays by materials is the photon–electron interaction. As it will be seen along this chapter, the complex amplitude (or modulus and phase) of the electromagnetic wavelets associated to photons elastically scattered in all directions
 by any material, is related to the tridimensional electron !

density function ρ r

through a Fourier transformation.

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Small-Angle X-ray Scattering by Nanostructured Materials

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 ! The electron density function ρ r fully describes the structure of materials; thus the ultimate goal of crystallographers and materials scientists is to determine this function, starting from experimental X-ray scattering patterns. Although this detailed information is not in practice fully obtained, relevant and useful structural features can generally be inferred. A typical SAXS setup is schematically shown in Fig. 1a. This technique provides useful structural information about heterogeneities in electron density sized within the range ~5 to 500 Å, these limits depending on the photon energy, sample-todetector distance, size of the beam-stopper and geometry of the X-ray detector. Very large objects as compared to the X-ray wavelength (with a size above, say, 1 μm) produce noticeable scattering intensity only within an extremely small angular

Fig. 1 (a) Schematic SAXS setup. (b) X-ray beam paths from the source (left) to the detector (right), both located far away from the sample. The total segment Δs ¼ AB þ BC is the optical path ! difference associated to the X-ray scattering by electrons in two elements of volume d r , from which the phase shift is determined

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domain close to the direction of the incident beam. Thus in this case the scattered photons hit the incident beam-stopper and are not recorded by the X-ray detector. Notice that the X-ray scattering intensity patterns within the “small-angle” range do not contain any information about the very short wavelength modulations in electron density associated to the atomic nature of the material, the effects from them only appearing in the scattering intensity profiles recorded at wide angles.

Basic Theory General Equations The intensity associated to electromagnetic waves elastically scattered by an electron was derived by Thompson. Since the amplitude of the X-ray wave scattered by an electron has a well-defined phase relation with the amplitude of the incident wave, interference between scattered wavelets occurs. For a nonpolarized incident X-ray beam with intensity I0, the intensity associated to the wavelets scattered by one electron per unit solid angle Ω, is I e ð2θÞ ¼ I 0 ½ð1 þ cos2 2θÞ=2 :r 2e , where 2θ is the scattering angle (i.e., the angle between the wave-vectors of incident and scattered photons), and re is the classical electron radius. The X-ray intensity associated to the elastic scattering by an electron at small angles per unit of incident beam intensity and per unit of solid angle can be considered as a constant, I e ¼ r 2e . Using this approximated value, the relative error in Ie for 2θ up to 8 is indicating the average over
the analyzed sample volume V. !

As indicated by Eq. 6 the function γ r

– named correlation function (Debye

and Bueche 1949) – is the volume average of the product of
Δρ in two volume ! ! elements connected by the vector r . The correlation function γ r is determined
 ! from the experimental scattering function I q by inverse Fourier transformation:
 ! γ r ¼

1 ð2π Þ3 V

ð
 ! ! ! ! I q ei q : r d q

(7)



 ! ! Provided that ρ r is known, the correlation function γ r can be determined

 ! ! by applying Eq. 6. But, inversely, from a known γ r function, ρ r cannot be unambiguously inferred. For isotropic systems,
 the correlation function is independent of the direction of !

!

the vector r , i.e., γ r becomes γ(r) and, consequently, the scattering intensity, I(q), ! !

i q  r is replaced in Eq. 5 by its is also isotropic. For D isotropic E systems the function e ! !

spherical average ei q : r

Ω

¼ sin qr=qr. Thus Eqs. 5 and 7, respectively, become

I ð qÞ ¼ V

ð1

4πr 2 γ ðr Þ

0

γ ðr Þ ¼

ð1

1 3

ð2π Þ V

0

sin qr dr qr

4πq2 I ðqÞ

sin qr dq qr

(8) (9)

A useful procedure that is often applied to characterize low-resolution structures, circumventing the phase problem,
is to  begin with an initial model described by a !

guessed electron density function ρ r . The scattering amplitude is thus determined

 ! ! by using Eq. 1 and then the resulting scattering intensity I q ¼ jA q j2 is

compared to the experimental intensity function. The use of ad hoc computer programs allows for many iterations and modifications of the structure model, until a good fit of the calculated function to the experimental curve is achieved. This procedure is, for example, applied to the determination of low-resolution structures (envelope functions) of proteins in dilute solution (Svergun 1999). For materials consisting of isolated (in general nonidentical) nanoparticles embedded in a homogeneous matrix, the scattering intensity I(q) is often modeled under the assumption of simple shapes and taking also into account eventual effects from spatial correlation. The model function is then fitted to the experimental scattering curves. An eventual good fitting justifies a posteriori the proposed model and yields the adjusted parameters that characterize the structure of the studied material.

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In another procedure, which is often applied to study structural transformations in materials subjected to isothermal annealing, the isotropic correlation function γ(r) is theoretically determined starting from basic thermodynamic and/or statistical concepts (Cahn 1965; Lebowitz et al. 1982). This is followed by the determination of I(q) for increasing periods of time using Eq. 8 and further comparison of the series of model functions with the sequence of experimental SAXS curves determined in situ along the structural transformation. This procedure is applied, for example, to verify the correctness of theoretical models of particle growth and structure coarsening.

Small-Angle Scattering by Nanoscopic Two-Phase Systems: Porod’s Law This section deals with isotropic biphasic materials, i.e., isotropic two-electron density systems with sharp interfaces, such as those schematically drawn in Fig. 2a, b. In this model the relevant parameters are the electron densities ρ1 and ρ2 and the volume fractions φ1 and φ2=1φ1. This model is applied to characterize different nanostructured materials such as nanoporous solids, nanocrystals, or disordered nanoclusters embedded in solid or liquid media, etc. The general properties of Fourier analysis tell us that the asymptotic trend, at high q, of the scattering intensity I(q) is connected to the behavior of the γ(r) function at small r. For isotropic two-electron density systems, the correlation function γ(r) can be approximated at small r by (Porod 1982):  γ ðr Þ ¼ ðρ1  ρ2 Þ φ1 ð1  φ1 Þ 1  2

S r 4Vφ1 ð1  φ1 Þ

 (10)

Fig. 2 Schematic examples of two types of biphasic structures or two-electron density systems. (a) Set of isolated spherical nano-objects with a constant electron density ρ1 embedded in a homogeneous matrix with electron density ρ2. (b) Bicontinuous structure, both phases with constant electron densities ρ1 and ρ2

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Small-Angle X-ray Scattering by Nanostructured Materials

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where S/V is the area of the interface per unit sample volume. Replacing γ(r) given by Eq. 10 in Eq. 8 and solving the integral, the leading term of the asymptotic intensity I(q), at high q, is given by (Porod 1982) I ð qÞ ¼

2π ðρ1  ρ2 Þ2 S q4

ð q ! 1Þ

(11)

Equation 11, named Porod law, applies to isotropic two-electron density systems with sharp interfaces, such as disordered porous materials and other two-phase systems whose relevant structural feature is their interface surface area. Porod’s law applies to either dilute or concentrated systems of isolated nanoobjects, provided they are not very thin sheets or very narrow cylinders, for which the asymptotic intensities are proportional to 1/q2 and to 1/q, respectively (Shull and Roess 1947). Equation 11 does not hold for sets of identical spherical or cylindrical nano-objects, because in these cases the SAXS intensity exhibits oscillations even at very high q. By analyzing the features of such oscillations, it is possible to determine the distance between the parallel portions of the interfaces (Ciccariello 1991). However, if the spherical or cylindrical nano-objects have a wide size distribution, the oscillations smear out and the asymptotic Porod’s law holds. For anisotropic two-electron density systems, Porod’s law still applies along ! all q directions, but the parameter S in Eq. 11 has a different meaning (Ciccariello et al. 2002). The behavior of I(q) at high q is often analyzed using I(q)q4 vs. q4 plots. Equation 11 implies that I(q)q4 becomes asymptotically constant in the high-q limit but, for many materials, the SAXS intensity also contains an additional and q-independent contribution from short-range density fluctuations in their phases (Ruland 1971). For these materials, the asymptotic I(q)q4 vs. q4 plot, at high q, is expected to exhibit a linear dependence (i.e., I(q)q4 = a + b q4) with a positive slope (b > 0). Extrapolation of the linear portion of the I(q)q4 function toward q4 = 0 yields I(q)q4(q = 0) = a. By substituting this value in Eq. 11, the interface area between both phases, S, is determined. On the other hand, Ruland (1971) demonstrated for two-phase systems with a smooth transition in electron density between both phases, that the asymptotic I(q)q4 vs. q4 plot at high q also exhibits linear dependence but in this case the slope is negative (b < 0). FromðEq. 9 it can be verified that γ(0) = Q/(2π 2V ), where Q is the integral defined as Q ¼

1

q2 :I ðqÞdq. On the other hand, for two-electron density systems, γ(0) is

0

equal to ðρ1  ρ2 Þ2 φ1 ð1  φ1 Þ (Eq. 10), so as the integral Q becomes` Q¼

ð1

q2 I ðqÞdq ¼2π 2 ðρ1  ρ2 Þ2 Vφ1 ð1  φ1 Þ

(12)

0

The integral Q depends on the electron density contrast factor (ρ1 – ρ2)2 and volume fractions of both phases but not on the specific features of their geometrical

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configuration. For example, along structural transformations that preserve both electron densities and phase volume fractions, even though the structure and, consequently, the shape of the scattering intensity curves change, the integral Q is expected to remain constant. Therefore, the integral Q is named “Porod invariant.” Examples of transformations that occur without significantly affecting the value of the integral Q are the processes of growth of homogeneous nanoclusters by mechanisms of coarsening or coalescence. For the determination of the interface surface area S by applying Eq. 11 the measurement of the scattering intensity in absolute units is required (See “Appendix: Experimental Issues”). Moreover, from Eqs. 11 and 12, the following equation is derived: ½I ðqÞq4 q!1 S ¼ π:φ1 ð1  φ1 Þ V Q

(13)

Thus, if the scattering intensity is only known in relative scale and provided the phase volume fractions are known, Eq. 13 allows for the determination of the specific interface surface area (S/V). Equation 13 is often applied to powdered samples, for which the precise measurement of the scattering intensity in absolute units is difficult.

Small-Angle Scattering by Spatially Uncorrelated Nanoparticles: Guinier’s Law The wavelets associated to the X-ray scattering by a dilute set of spatially uncorrelated nano-objects do not interfere. Under this condition and provided the objects are identical and centrosymmetric, the total scattering intensity I(q) is expressed as I ðqÞ ¼ NI 1 ðqÞ

(14)

where N is the number of nanoparticles and I1(q) is the SAXS intensity produced by a single nanoparticle. By solving Eq. 8 for an arbitrary correlation function γ(r) associated to a single nano-object, it can be demonstrated (Guinier and Fournet 1955) that the SAXS intensity at small q is given by I ðqÞ ¼ N ðΔnÞ2 eRg q

2 2

=3

ð q ! 0Þ

(15)

where Δn is the excess in number of electrons inside the nano-objects and Rg their radius of gyration. For nano-objects with volume V1 and constant electron density ρ1, embedded in a homogeneous matrix with electron density ρ2, the number of electrons in excess is Δ n ¼ ðρ1  ρ2 ÞV 1 and the radius of gyration is

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Small-Angle X-ray Scattering by Nanostructured Materials

ð Rg ¼

!

1195

1=2

r 2 d r =V 1

(16)

V1

Equation 15 is named Guinier law. In order to derive the radius of gyration Rg of nano-objects from results of SAXS measurements, the Guinier plot (log I vs. q2) is applied. In this plot, a straight line is expected to be observed at small q, within a more or less wide q range depending on the size and shape of the objects (Guinier and Fournet 1955). From the slope αG of the straight line in Guinier plots, the radius of gyration is determined; Rg ¼ ½3ðαG =log:eÞ1=2 ¼ 2:628:jαG j1=2 . For example, the radius of gyration of homogeneous spherical objects is related to their radius R by Rg = (3/5)1/2R and that of homogeneous cylinders with radius     1=2 R and height H by Rg ¼ R2 =2 þ H 2 =12 . Guinier plots are also applied to determine the SAXS intensity at q = 0, I(0), by linear extrapolation of log I (q2) to q2 = 0. The total SAXS intensity produced by a dilute set of nano-objects with a distribution of radii of gyration N(Rg) is given by the sum of the individual contributions of each object. For this system Guinier’s law also holds but the derived parameters are weighted averages. For example, for two-electron density systems consisting of a isotropic and polydisperse set of N spatially uncorrelated nanoobjects, Eq. 15 becomes

 2 2 I ðqÞ ¼ N ðρ1  ρ2 Þ2 V 21 ehRgiG q =3 ðq ! 0Þ

(17)

where is the average of V21 and < Rg >G is a weighted average (named Guinier average) defined as 2ð

31=2   2 2 V N R R dR g g7 1 g 6

 7 Rg G ¼ 6 4 ð   2 5 N Rg V 1 dRg

(18)

ð

  with N Rg dR ¼ N. Notice that the G averaging weights more large objects than small ones. For a polydisperse set of spherical nano-objects, Eq. 18 becomes 2ð

31=2 " #1=2 8 N ð R ÞR dR 8 6 7 7 ¼ R  hRiG ¼ 6 4ð 5 R6 N ðRÞR6 dR

(19)

Guinier law is usually applied to determine the radius of gyration of nano-objects with narrow size distribution. For highly polydisperse systems, the q range over

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which Guinier law holds is small and Guinier plot yields a weighted average of the radius of gyration far from the arithmetic average and strongly biased toward those of the largest objects. This effect is schematically described for two sets of spherical objects with same arithmetic average radius = 40 Å but different widths of radius distribution, as shown in Fig. 3a. For these two systems the slopes of the linear portion of Guinier plots and, consequently, the average radius < R>G derived by applying Guinier law are different (Fig. 3b). The extrapolated intensity I(0) for polydisperse systems, being proportional to the average , also depends on the shape of the radius distribution. From Eq. 15, it can be inferred that the experimental SAXS intensity extrapolated to q = 0 corresponding to a dilute set of N identical objects is given by I(0) = N (Δn)2. For two-electron density systems composed of a dilute set of N nano-objects, each of them with volume V1 and electron density ρ1, embedded in a matrix with electron density ρ2, the SAXS intensity at q = 0 is I ð0Þ ¼ N ðρ1  ρ2 Þ2 V 21

Fig. 3 (a) Narrow N1(R) and wide N2(R) radius distributions of spheres with same arithmetic radius average = 40 Å. (b) Guinier plots of SAXS intensities G1 and G2, at small q, corresponding to the radius distributions N1(R) and N2(R), respectively. The magnitude of the slope of the linear part of the log I vs. q2 plot at small q and the extrapolated intensity I (0) associated to the radius distribution N2(R) are both higher than for N1(R)

(20)

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Small-Angle X-ray Scattering by Nanostructured Materials

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For dilute solutions we have φ1 V ¼ NV 1 and ð1  φ1 Þ  1 so as Eq. 12 becomes Q ¼ 2π 2 N ðρ1  ρ2 Þ2 V 1 . Thus, regardless the object shape, its volume V1 can be determined from the quotient I(0)/Q as follows: V 1 ¼ 2π 2

I ð 0Þ Q

(21)

Equation 21 can also be applied to polydisperse systems, the result being in this case the quotient between averages, /. Identical anisotropic objects with same orientation produce anisotropic scattering ! patterns, i.e., the scattering intensity depends on the direction of the vector q . In the ! limit of small q , Guinier law becomes (Guinier and Fournet 1955) I ðqD Þ ¼ N ðρ1  ρ2 Þ2 V 21 eRD qD 2

2

(22)

!

where qD refers to the component of q in the direction along which the scattering intensity is measured and RD is the inertia distance of the object in the same direction, from a perpendicular plane containing the center of “mass” of the electron density function. If the system is composed of identical and anisotropic nano-objects that are randomly oriented, the resulting scattering intensity is isotropic. In this case, the structural parameter determined by applying Guinier law (Eq. 15) is the radius of gyration of the nano-objects.

Dilute Sets of Nanoparticles Spherical Nanoparticles Schematic views of monodisperse and polydisperse sets of spherical nano-objects are shown in Fig. 4a, b, respectively. The scattering intensity associated to a single spherical and homogeneous nano-object embedded in a homogeneous matrix, with spatially constant electron densities ρ1 and ρ2, respectively, is derived from the amplitude A1(q) defined by Eq. 3. For a spherical nano-object with radius R the scattering intensity is given by  2 ðR sin qr dr I 1 ðqÞ ¼ jA1 ðqÞj2 ¼ ðρ1  ρ2 Þ 4πr 2 qr 0

(23)

By solving the integral, Eq. 23 becomes  2 I 1 ðqÞ ¼ ðρ1  ρ2 Þð4π=3ÞR3 Φðq, RÞ where Φ(q, R) is

(24)

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Fig. 4 Schematic views of systems composed of dilute sets of (a) monodisperse and (b) polydisperse sets of spherical nano-objects. (c) Scattering intensities corresponding to three samples containing spherical objects with the same average radius = 40 Å and a Gaussian distribution N(R) with three different standard deviations: σ = 0, σ = 5 and σ = 15 Å

Φðq, RÞ ¼ 3

sin qR  qR cos qR ðqRÞ3

Thus the total scattering intensity produced by a dilute (spatially uncorrelated) set of N identical spheres is I(q) = NI1(q), i.e., I ðqÞ ¼ N ðρ1  ρ2 Þ2



4π 3: 2 R ½Φðq, RÞ2 3

(25)

The scattering intensity given by Eq. 25 is plotted in Fig. 4a, for identical spheres with radius R = 40 Å. At high q the intensity function exhibits several secondary maxima and zeros, the different zeros being located at qR = 4.50, 7.72, 10.90 . . . The scattering intensity related to a dilute set of N spherical nano-objects with a radius distribution defined by N(R) is calculated by ð I ðqÞ ¼ N ðRÞI 1 ðq, RÞdR

(26)

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Small-Angle X-ray Scattering by Nanostructured Materials

1199

where I1(q, R) is the scattering intensity produced by a single sphere (Eq. 24). The scattering intensity curves related to three dilute sets of spherical objects, with different Gaussian radius distributions and same arithmetic average radius = 40 Å, are plotted in Fig. 4c. The standard deviations of the Gaussians are σ = 0 (monodisperse system), 5 Å and 15 Å. It can be noticed in Fig. 4 that, for increasing polydispersivity, the secondary maxima and zeros progressively smear out. On the other hand, in this case the intensity I(0), being proportional to or , is higher for wider radius distributions. The radius distribution N(R) of a dilute and polydisperse set of spherical nanoparticles can be derived from the measured I(q) functions by solving the integral Eq. 26. For this purpose, the program package named GNOM (Svergun 1992) is often used. The output of GNOM yields the volume weighted distribution function, D(R), related to N(R) for spheres by D(R) = (4π/3)R3N(R). GNOM is also applied to determine the volume distribution function of nano-objects with other simple shapes. Moreover the intensity function I1(q) related to objects with complex shapes can be independently determined and used as an input file in GNOM program.

Application (Example 1): PbTe Nanocrystals Embedded in a Silicate Glass An experimental SAXS study of a system composed of PbTe nanocrystals embedded in a silicate glass was performed by Craievich et al. (1997). This nanostructured material exhibits interesting nonlinear optical properties in the infrared, making it potentially useful for applications to telecommunication devices. A silicate glass doped with Pb and Te was held at high temperature, quenched by splat-cooling down to room temperature and then submitted to an isothermal annealing at 650  C. Initially, isolated Pb and Te atomic species diffuse through the supersaturated glass and nucleate PbTe nanocrystals which progressively grow. A number of SAXS intensity curves were successively recorded in situ, along the whole annealing process. The experimental results are displayed in Fig. 5. The SAXS intensity progressively increases for increasing annealing time. At high q, the intensity curves exhibit satellite peaks or secondary maxima that are characteristic of the scattering function associated to a set of spheres with nearly identical radius. The secondary maxima progressively shift toward smaller q, as expected for a set of growing nanospheres (Eq. 25). Because of the high statistical dispersion in the scattering intensities at high q, the secondary maxima are not clearly apparent in the curves corresponding to early stages of nanocrystal growth. The positive deviation of the experimental points from the theoretical modeled curve, at very small q, indicates the existence of additional and rather large heterogeneities in electron density in the glass matrix. The experimental SAXS curves displayed in Fig. 5 were well fitted by model functions defined by Eq. 26, which applies to dilute sets of spherical objects, assuming a time-varying average nanocrystal radius and a Gaussian radius distribution, N(R), with a time-independent relative standard deviation σ/ = 0.08.

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Log intensity (arbitrary units)

10

8

6

4

2

0 –0.2

–0.1

0.1

0.0

0.2

0.3

q (Å–1)

Fig. 5 Scattering intensity curves recorded in situ, corresponding to a dilute set of spherical PbTe nanocrystals embedded in a homogeneous silicate glass, during isothermal growth at T = 650  C. The period of time for nanocrystal growth increases from 19 up to 119 min from bottom to top. The continuous line is the best fits of Eq. 26 using a Gaussian N(R) function with a time-varying radius average and a constant relative standard deviation σ/ = 0.08. The curves are vertically displaced for clarity (Reprinted with permission from Craievich et al. (1997). Copyright 1997 by the International Union of Crystallography)

For the sample held 2 h at 650  C, the best fit of the model scattering curve led to = 32.5 Å and σ = 2.6 Å. The time dependence of the average radius agrees with the prediction of the classical theory for nucleation and growth of spherical precipitates in a homogeneous matrix.

Application (Example 2): Clustering of Colloidal ZnO Nanoparticles Powders consisting of ZnO nanoparticles produced by sol-gel route are used as precursors for developments of new materials with interesting properties. The first step of the sol-gel route leading to ZnO solid nanoparticles is the formation of a liquid suspension of zinc acetate in ethanol, to which LiOH is added under ultrasound treatment. An in situ SAXS study was performed in order to characterize the first steps of aggregation of ZnO nanoparticles in liquid solution (Tokumoto et al. 1999). The different experimental scattering functions, recorded after increasing periods of time at 40  C, were analyzed by assuming that the system is dilute and that the colloidal nano-objects are spherical. In order to determine the radius distribution of the particles, the integral Eq. 26 was solved by using GNOM program (Svergun and Semenyuk 1991; Svergun 1992). GNOM was applied to all experimental scattering curves of the studied ZnO-based suspension corresponding to different aggregation

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Small-Angle X-ray Scattering by Nanostructured Materials

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e Tim

Fig. 6 Time-dependent volume weighted radius distribution, D(R), for ZnO-based colloidal particles in liquid suspension maintained inside a sealed cell during SAXS measurements. The time of growth increases from 10 up to 120 min. The D(R) functions were derived from the set of experimental SAXS curves by applying the GNOM program (Reprinted with permission from Tokumoto et al.(1999). Copyright 1999 by Elsevier)

150

100

50

0

R (Å)

times, thus yielding the set of volume weighted radius distribution functions D(R) plotted in Fig. 6. The shape of the D(R) function and its time variation (Fig. 6) suggested that the kinetics of formation of ZnO clusters is characterized by two main stages. During the first stage, a growing peak centered at R = 17 Å is apparent, indicating a continuous formation of small clusters. The number of clusters increases monotonously for increasing reaction time, while their average radius, = 17 Å, remains constant. During the second stage, the volume weighted distribution exhibits a still growing peak at 17 Å, while the formation and growth of a second peak corresponding to an initial average particle radius = 60 Å is also apparent. This peak shifts continuously toward higher R values, up to 110 Å, along a period of time of 2 h. The described time variation of the volume weighted distribution function clearly evidences the continuous formation of colloidal primary clusters and their simultaneous aggregation and growth.

Concentrated Sets of Nanoparticles Spatially Correlated Spherical Nanoparticles Many sol-gel based isotropic nanomaterials consist of spatially correlated nanoparticles embedded in a homogeneous matrix. Examples are concentrated colloidal sols (solid nanoclusters embedded in a liquid medium) and solid hybrid

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nanomaterials (inorganic clusters embedded in a solid polymeric matrix). Two models of SAXS functions associated to different types of systems composed of spatially correlated nano-objects will be described, one of them containing identical nanoclusters and another consisting of a two-level hierarchical structure. The total scattering intensity produced by a set of identical and spatially correlated nano-objects is affected by interference effects, thus Eq. 14 does not hold. For isotropic systems composed of a set of N spatially correlated spherical (or more generally centrosymmetrical) nano-objects, the SAXS intensity is given by I ðqÞ ¼ NI 1 ðqÞSðqÞ

(27)

where S(q) is the structure function that accounts for interference effects produced by spatial correlation. For a set of nano-objects without long-range order, the structure function S(q) tends asymptotically to 1 at high q. For a set of spatially uncorrelated nano-objects SðqÞ ¼ 1 over the whole q domain, and thus Eq. 27 becomes equivalent to Eq. 14. A semiempirical structure function that is often applied to describe spatial correlation in isotropic systems composed of spherical nano-objects embedded in a homogeneous matrix, derived using the Born–Green approximation, is given by (Guinier and Fournet 1955): Sð qÞ ¼

1 1 þ kΦS ðqÞ

(28)

where k, named “packing factor,” is associated to the degree of compactness of the local structure (for the closest packing of spheres kmax is equal to 5.92) and Φs(q) is Φ S ð qÞ ¼ 3

sin qd  qd cos qd ðqdÞ3

(29)

where d is the average distance between the spatially correlated nano-objects. Several examples of models for scattering intensity functions are displayed in Fig. 7a, b. These functions are determined by Eq. 27 with S(q) given by Eq. 28 for a set of spheres with same radius, R = 10 Å, and different d and k values. The intensity curves displayed in Fig. 7a show that the q value corresponding to the maximum of the scattering curves, qmax, decreases for increasing average distances. On the other hand, the different curves plotted in Fig. 7b indicate that increasing values of packing factor k yield more pronounced and well-defined scattering peaks. A rough estimate of the average distance between particles is usually inferred by applying the simple equation d = 2π/qmax. However, by analyzing the curves plotted in Fig. 7a, it can be verified that the equation d = 5.6/qmax yields a better estimate of the average distance. Anyway, the determination of a more precise average distance between particles requires the fitting of a model intensity function to the whole experimental intensity curve. Even for nano-objects that are not spherical but instead exhibit a globular shape, the structure function given by Eq. 28 is usually applied as a good approximation.

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Small-Angle X-ray Scattering by Nanostructured Materials

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Fig. 7 Model scattering intensity curves corresponding to different sets of spatially correlated spheres, all of them with same radius, R = 10 Å. (a) Packing factor k = 3 and average interparticle distances: d = 30 Å, d = 50 Å, and d = 70 Å. (b) Average distance d = 50 Å and packing factors k = 1, k = 3 and k = 5. The normalized scattering intensity curve for a dilute set of particles with same radius is displayed as a black line in (a) and (b)

This structure function is also applied to model scattering intensity curves associated to materials composed of polydispersed nano-objects with narrow radius distributions. Many hybrid materials prepared by the sol-gel process were studied by SAXS. Some of these hybrid materials are composed of a isotropic set of inorganic nanoclusters embedded in a polymeric matrix. The heterogeneous nature of these nanostructured materials is characterized by using a simple two-electron density model consisting of high electron density clusters embedded in a low electron density matrix (Dahmouche et al. 1999). Certainly, the polymeric phase exhibits electron density fluctuations at molecular level that also produces small-angle scattering, but their contribution to the total scattering intensity is assumed to be weak and/or not strongly varying with q. The basic assumption here is that the dominant contribution to small-angle scattering intensity comes from the electron density contrast between inorganic nanoclusters and polymeric matrix. Some materials are heterogeneous at multiple scale levels. For example, nanometric clusters may segregate and form cluster-rich domains embedded in a

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cluster-depleted matrix. For this particular two-level system, the effects on the SAXS intensity produced by a coarse structural level and another fine level are expected to be dominant at low and high q, respectively. For the example to be described in the next section corresponding to a isotropic two-level structure – with its fine level consisting of spatially correlated nano-objects – the scattering intensity can be modeled by the following semiempirical equation (Beaucage et al. 1995):   n   3 oP1 2 2 2 2 I ðqÞ ¼ G1  eð1=3ÞRg1 q þ B1  eð1=3ÞRc q erf qRg1 =61=2 =q   n   oP2 ð1=3ÞR2g2 q2 1=2 3 þ G2  e þ B2  erf qRg2 =6 =q  Sð qÞ

(30)

where sub-indexes 1 and 2 refer to the coarse and fine structure levels, respectively. The factors Gi are equal to Ni(Δni)2 (Eq. 15) and Bi are related to Gi by specific equations that depend on the object geometry, and Pi are Porod exponents that are equal to 4 for simple two-electron density systems and may have other values depending on the geometry of the objects. The second term in Eq. 30 corresponding to the fine level also includes the structure function S(q) accounting for spatial correlation of the small clusters inside the volume defining the coarse level. In the first associated to the coarse structure, the Gaussian function  term,  2 2 given by exp Rc q =3 is a high-q cutoff factor in which Rc = Rg2 (Beaucage et al. 1995). Provided the X-ray scattering experiment covers a wide q range, hierarchical structures consisting of more than two structure levels can also be characterized. In order to model SAXS intensity curves associated to these complex materials, additional terms are included in Eq. 30. Since the q range to be covered for the study of many-level structures is rather wide, several SAXS measurements with the same sample but using different collimation conditions, sample-to-detector distances and/or X-ray wavelengths, are required. Examples of fittings of model functions assuming multilevel structures to a number of experimental SAXS curves were reported by Beaucage et al. (1995). In order to characterize coarse structures composed of very large (micrometric) particles, the use of ultra-small-angle X-ray scattering (USAXS) – q range below 0.001 Å1 – or light scattering techniques is required.

Application: Fe-Doped Organic–Inorganic Hybrid Nanomaterials Many organic–inorganic composite materials exhibit interesting properties that can be tailored by an adequate control of the preparation conditions (Dahmouche et al. 1999). Moreover, the structural characterization of these materials is needed in order to explain their magnetic behavior. The structure of a number of hybrids composed of spatially correlated siliceous nanoparticles or clusters embedded in a

39

Small-Angle X-ray Scattering by Nanostructured Materials

1205

a

Intensity (a.u.)

10

1

0.1 0.02

0.04

0.06 0.08 0.1

0.2

0.4

−1

q (A )

b Level 2 Level 1

Fig. 8 (a) Experimental scattering intensity produced by siliceous clusters containing 0.76 wt% Fe(II) embedded in a polymeric matrix. The continuous line is the best fit of Eq. 30 to the experimental curve. The dashed lines indicate the Guinier and Porod contributions to the scattering intensity produced by siliceous clusters and the structure function (oscillatory curve). The dotted lines are the Guinier and Porod contributions to the scattering intensity associated to the coarse domains. (b) Schematic view of the proposed two-level model. The small circles correspond to siliceous clusters (Figure 8a reprinted with permission from Silva et al. (2003). Copyright 2003 by the International Union of Crystallography)

matrix consisting of grafted polymeric chains were well described by a two-electron density model. For these systems the SAXS patterns exhibit a correlation peak located at decreasing q values for increasing molecular weight of the polymer molecule (Dahmouche et al. 1999). A SAXS study of hybrid organic–inorganic nanomaterials composed of Fe(II)doped di-ureasils was carried out by Silva et al. (2003). Figure 8a displays the scattering intensity produced by a di-ureasil hybrid doped with 0.76 wt% Fe(II). In

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A. F. Craievich

order to characterize the structure of Fe(II)-doped nanohybrids, the two-level model described in the precedent section (Beaucage et al. 1995) was applied. The SAXS intensity corresponding to the fine structure level displays a peak associated to cluster-cluster correlations, centered at q = 0.15 Å1, which is also observed for undoped samples. For Fe(II)-doped hybrids, this peak is slightly shifted toward higher q. For q < 0.1 Å1, the scattering intensity is mainly related to the coarse structural level. The model scattering curve defined by Eq. 30 for two structural levels, including the structure function S(q) for the fine level given by Eq. 28, is displayed in Fig. 8a. This figure also shows the Guinier and Porod contributions to the total scattering intensity corresponding to both levels. The radii of gyration Rg obtained by the best fit procedure are 7.5 Å for the small clusters and 54 Å for the coarse domains. Similar analyses of SAXS curves for different Fe(II) doping levels, up to 4.5 wt%, revealed a decreasing average distance between siliceous clusters for increasing Fe(II) content. This result suggests that Fe(II) ions are dispersed in the polymeric matrix, these ions promoting a shrinkage effect that leads to the observed decrease in average cluster-cluster distance. The model structure consisting of large domains containing spatially correlated siliceous particles embedded in a depleted matrix is schematically shown in Fig. 8b. The reported results indicate that the formation of coarse silicide-rich domains is promoted by the addition of Fe(II) ions (Silva et al. 2003).

Fractal Structures Small-Angle Scattering by Fractal Structures The SAXS method is applied to structural characterization of a number of materials which exhibit a self-similar or fractal structure, and also to the determination of the mechanisms involved in aggregation processes, either in precursor sols or after the sol-gel transition. Fractal materials are characterized by three relevant structural parameters: (i) a radius r0 corresponding to the size of the individual primary particles (basic nanoobjects that build up the fractal structure), (ii) a fractal dimension D that depends on the nature of the mechanism of aggregation, and (iii) a correlation length ξ that defines the size of isolated aggregates or the cutoff distance of the fractal structure for percolated systems such as fractal gels. A homogeneous object and another with fractal structure – built up by N small primary particles – are schematically shown in Fig. 9a, b, respectively. The number of primary particles inside a sphere of radius r, measured from the center of mass of the fractal aggregate, is given by  D N ðr Þ ¼ r=r 0

(31)

39

Small-Angle X-ray Scattering by Nanostructured Materials

1207

Fig. 9 (a) Homogeneous object and (b) fractal object. The red curve in (c) is the scattering intensity, I(q) = I1(q).S(q), associated to a fractal object with size of primary building blocks ro = 5 Å, correlation length ξ = 5000 Å, and fractal dimension D = 1.80. The scattering intensity corresponding to the primary particles, I1(q) (olive), and the structure function, S(q) (blue), are also plotted in (c). Notice that for large fractal aggregates the determination of the correlation length requires SAXS measurements down to very low minimum q value (qmin  1/ξ)

where r0 is of order of the size of the primary basic units that build up the fractal object. Thus the mass M(r) inside a sphere with radius r, for both (homogeneous and fractal) objects, is proportional to rD, the exponent being D = 3 for homogeneous objects and D < 3 for fractal aggregates. The SAXS intensity associated to a correlated set of primary nanoparticles building up a fractal structure is defined by Eq. 27, which involves the scattering intensity by single primary particles, I1(q), and the structure function, S(q), associated to the nature of their spatial correlation.

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A. F. Craievich

Different simple functions have been used for I1(q), such as the intensity produced by spherical particles (Eq. 25) or the Debye–Bueche function, defined by I 1 ð qÞ ¼ 

A 1 þ r 20 q2

2

(32)

where A is a constant. The structure function S(q) corresponding to a fractal object is derived from the radial distribution function for primary particles inferred from Eq. 31. This distribution function is multiplied by a cutting function that defines a structural correlation length ξ. This analysis finally leads to the following structure function (Teixeira 1988): Sð qÞ ¼ 1 þ

1 ðqr 0 Þ

D

h

 D  Γ ðD  1 Þ 1 iðD1Þ=2 sin ðD  1Þ tan ðqξÞ 1 þ 1=ðqξÞ2

(33)

where Γ is the gamma function. Thus, by selecting I1(q) defined by Eq. 32 and the structure function S(q) given by Eq. 33, the scattering intensity produced by a fractal aggregate, or by a set of spatially uncorrelated fractal aggregates, is I ð qÞ / 

1

2 1 þ r 20 q2 9 8 > > <  = 1 D  Γ ð D  1Þ 1  1þ sin ðD  1Þ tan ðqξÞ h i > > ðqr 0 ÞD 1 þ 1=ðqξÞ2 ðD1Þ=2 ; :

(34)

A scattering intensity function defined by Eq. 34 for particular values of the three structural parameters (ro, ξ, D) is plotted in log–log scale in Fig. 9c. Since the size of the primary particles is much smaller than the correlation length, I1(q) is constant within a rather wide low-q range, thus the variation of the scattering intensity at small q’s is dominated by the structure function. At high q, S(q) becomes a constant (S(q) = 1) and thus the variation in the scattering intensity in this q range is governed by I1(q). Figure 9c displays a log I vs. log q plot associated to a fractal object with correlation length much larger than the size of the primary particles (ξ r0). We notice in this log–log plot the presence of three q ranges over which linear dependences with different slopes are apparent: (i) Over the small q range (q  l/ξ) the slope is zero. In this q range the scattering intensity behaves as expected from Guinier’s law, its value extrapolated to q = 0, I(0), being related to the fractal dimension D by

39

Small-Angle X-ray Scattering by Nanostructured Materials

I ð 0Þ / ξ D

or

I ð0Þ / RD g

1209

(35)

with Rg ¼ f2=½DðD þ 1Þg1=2 ξ: (ii) Over the intermediate q range, i.e., for 1/ξ  q  1/r0, the magnitude of the slope is equal to the fractal dimension D. This implies that the scattering intensity exhibits a simple power q-dependence, I(q) / q–D. (iii) Over the high q range (q 1/r0) the slope is 4, this implying that Porod’s law (I(q)/ q4) holds. Two crossovers of the different linear parts in log I vs. log q plots, at q = q1 and q = q2 (q2 > q1), are shown in Fig. 9c. The radius of the primary particles r0 is simply related to q2 by r0 = 1/q2 and the size parameter of the fractal aggregate or correlation length is given by ξ = 1/q1. Thus, if ξ r0, the relevant structure parameters ξ, and r0 can be directly determined from log–log plots of the scattering intensity. If the condition ξ r0 is not satisfied no well-defined crossovers are apparent. In this case, the parameters ξ, D, and r0, are determined by fitting the I(q) function defined by Eq. 34 to the whole experimental curve. The fractal dimension D can also be determined by applying Eq. 35 to a set of experimental SAXS curves determined in situ, during an aggregation process. The values of I(0) and Rg are determined from Guinier plots (Log I(q) vs. q2) for all successive SAXS curves. Since I ð0Þ / RD G the plot of log I(0)–log Rg is expected to be linear, the slope of the straight line yielding the fractal dimension D. If the condition ξ r0 is not fulfilled the “fractal” model cannot be safely applied. It is a general consensus that, in order to establish the fractal nature of an aggregate, the quotient ξ/r0 should be of the order of or larger than 10. In addition, it must be remembered that power q-dependences leading to D values smaller than 3 are also expected for nonfractal objects such as, for example, narrow linear chains or thin platelets. Therefore, independent evidences supporting the use of fractal models are often required. Many mechanisms involved in aggregation processes were analyzed and the respective fractal dimensions of the resulting structures were theoretically determined (Meakin 1986). By associating these theoretical results with experimental determinations of the dimension D, the mechanisms that govern aggregation processes leading to fractal structures can be established.

Applications: Aggregation in Zirconia-Based Sols and Gels The formation of zirconia-based gels promoted by the aggregation of colloidal particles in sol state was investigated in situ by SAXS (Lecomte et al. 2000). All experimental scattering curves, plotted as log I(q) vs. log q in Fig. 10, exhibit a wide q range with well-defined linear behavior. Following the procedure described in the precedent section, the magnitude of the slope of the straight line was assigned to the

1210

I(q) (arb. units) 2 000

1 000

500

−1.7

3 000 1 000 INTENSITY (a.u)

Fig. 10 Log I vs. log q plots corresponding to a zirconiabased sol held at room temperature for increasing periods of time, from 4 h (bottom) up to 742 h (top). The inset is the scattering intensity curve of the final gel obtained after a period of about twice the gelling time (Reprinted with permission from Lecomte et al. (2000). Copyright 2000 by the International Union of Crystallography)

A. F. Craievich

300 100

−4

30 10 0.1

200

0.5 1 2 3 q (nm –1)

100

50

20 0.1

0.5

1

2

3

q (nm –1)

fractal dimension of the growing aggregates, D being equal to 1.7 along the whole aggregation process. The low-q limit of the linear portion of the scattering curves displayed in Fig. 10, and thus the crossover q1, progressively shifts toward lower q for increasing periods of time. This indicates that the aggregate size (ξ = 1/q1) continuously grows. The crossover q2 is not visible in the main set of curves displayed in Fig. 10 but, in the inset, corresponding to a SAXS curve determined up to a higher q value, this crossover toward a Porod behavior (I(q) / q4) is apparent. This suggests that the primary subunits have a smooth and well-defined external surface. The results reported by Lecomte et al. (2000) indicate that the fractal clusters in the studied zirconia-based sols are formed by aggregation of very small colloidal particles already existing at the beginning of the hydrolysis and condensation reactions. On the other hand, the maximum observed in the scattering curves for q 6¼ 0 is related to the existence of spatial correlations between the fractal aggregates, which could analytically be described by an inter-aggregate structure function S0 (q) defined in the same way as S(q), by Eq. 28, and included as another factor in Eq. 34. A fractal dimension close to that experimentally determined (D = 1.7) has been derived by computer simulation (Meakin 1986) for the mechanism of growth named diffusion-limited cluster-cluster aggregation (DLCA). Since the slope of all scattering curves displayed in log–log scale does not exhibit any variation with time, it

39

Small-Angle X-ray Scattering by Nanostructured Materials

1211

a

b

I(q) (arb. units)

I(0) (arb. units)

slope 1.78±0.06

0.1

c

b a

0.01 0.1

slope 0.98±0.06

1E−3 1

q (nm−1)

1

10 Rg(nm)

Fig. 11 Scattering intensity curves from sulfate-zirconia sols with different HNO3, H2SO4, and H2O contents. (a) Log–log plots of the scattering intensity produced by a few selected samples maintained inside a sealed cell at the end of their aggregation process. (b) Plot of I(0) vs. Rg, in log–log scale, corresponding to the final states of a number of sols with different compositions (Reprinted with permission from (Riello et al. 2003). Copyright 2003 by the American Chemical Society)

could be concluded that the fractal dimension D and, consequently, the mechanism of aggregation remains invariant during the whole aggregation process. Another SAXS study of sulfate-zirconia sols with several compositions (varying HNO3, H2O and H2SO4 contents) was reported by Riello et al. (2003). In order to characterize the aggregation mechanism, these authors determined successive SAXS curves after progressively increasing time periods keeping the sols in open cells. The values I(0) and Rg – determined by applying Guinier law (Eq. 15) to every scattering curve – were plotted as log I(0) vs. log Rg. This plot was analyzed by applying Eq. 35, which predicts for fractal objects a linear behavior with a slope equal to the fractal dimension D of the growing aggregates. The process of cluster growth in sulfate-zirconia sols with different compositions in sealed cells was also studied. Since, under sealed condition, the reactions in sols are very fast, only the scattering curves corresponding to the final states could be determined (Fig. 11a). The log I(0) vs. log Rg plot corresponding to the final states of all studied samples is displayed in Fig. 11b. Notice that the experimental points lie on two different straight lines, each of them with a slope similar to those observed by the same authors in previous in situ studies during the cluster growth in open cells. The slope of the straight line for Rg < 20 Å in the log I vs. log Rg plot displayed in Fig. 11b is close to D = 1.0 thus suggesting that the aggregation process starts by the

1212

A. F. Craievich

formation of short 1D linear chains. This initial regime is followed by another one involving the cross-linking of the precursor linear chains which build up a threedimensional fractal structure. The fractal dimension experimentally determined for Rg > 20 Å is 1.8, which is close to the expected theoretical value for diffusion-limited cluster-cluster aggregation. It was then concluded that, even though the sizes of the final aggregates in a number of sols, with very different compositions, vary from 0.5 nm up to 10 nm, the mechanism of growth of all of them is essentially the same. In the SAXS study reported by Riello et al. (2003) the mechanism of growth of the aggregates is theoretically characterized by an exponent D from the early stages of the clustering process, when the aggregates are still rather small and the condition ξ r0 for a fractal object is not yet fulfilled. Therefore, in these early stages, the exponent D – derived from in situ SAXS experiments by applying Eq. 35 – should not be assigned to a fractal dimension, but instead it must be considered as a useful parameter that characterizes the mechanism of growth.

Nanophase Separation General Considerations A number of nanoheterogeneous materials are formed by phase separation processes starting from a homogeneous solid solution at high temperature brought by fast cooling into a miscibility gap. In supersaturated and initially homogeneous (quenched) solid solutions with a composition close to the binodal curve (which defines the solubility limits), phase separation occurs by nucleation and growth of a minor new phase. This leads to a final two-phase material consisting of isolated and initially nanoscopic particles embedded in a homogeneous matrix (Fig. 2a). The growth of the second-phase particles can be characterized by in situ SAXS, using in this case a model consisting of a dilute or concentrated set of spherical particles surrounded by a solute-depleted shell. On the other hand, the final structure – after long periods of heat treatment – of initially homogeneous solid solutions brought, by quenching, close to the central part of a miscibility gap is described by a two-phase bicontinuous model, both phases occupying nearly the same volume fraction (Fig. 2b). For the first stages of phase separation occurring near the central part of the miscibility gap, a theoretical model named spinodal decomposition was proposed by Cahn (1965). At advanced stages of phase separation, even after having reached the equilibrium concentrations, both phases still exhibit a structural evolution driven by a pure coarsening mechanism.

Phase Separation and Dynamical Scaling Property In order to describe the advanced stages of nanophase separation (i.e., the coarsening regime) in binary materials, a statistical model was proposed by

39

Small-Angle X-ray Scattering by Nanostructured Materials

1213

Marro et al. (1975) and Lebowitz et al. (1982). This model assumes that the material contains atoms A and B arranged in a simple cubic lattice with an occupation function η(ri), which takes values +1 or 1 for sites ri occupied by atoms A or B, respectively. A probability function for atom exchanges and a simple equation for the energy of the system was proposed. This model is analogous to that applied to ferromagnetic Ising spin systems. Finally, the theoretical isotropic and time-dependent structure function, S(q, t), was determined by computer simulation. In the proposed model the primary particles are spatially correlated atoms whose scattering intensity I1(q) at small q is constant. Consequently, the SAXS intensity (Eq. 27) can be written as I ðq, tÞ / Sðq, tÞ

(36)

Different moments Sn(t) and normalized moments qn(t) of the structure function, S(q,t), are defined as Sn ð t Þ ¼

ð1 ð01

qn ð t Þ ¼

ð0

Sðq, tÞqn dq Sðq, tÞqn dq

1

(37)

Sðq, tÞdq

0

Marro et al. (1975) determined the time variation of the structure function S(q,t) and its associated moments at advanced stages of phase separation, after both phases having reached their final compositions. Their results of computer simulations demonstrated that the structure function and its moments exhibit the following properties: (i) The second moment remains invariant, S2(t) = S2. Since S2 is proportional to the integral Q (Eq. 12), its time invariance implies that the advanced stage of phase separation is governed by a pure coarsening process. (ii) The time variation of the structure function S(q,t) exhibits a dynamical scaling property, evidenced by the existence of a time-independent function F(x) given by Fð x Þ ¼

Sðq, tÞ ½q1 ðtÞds S2

(38)

where the coordinate x is equal to (q/q1) and ds is the dimension of the space in which the process of phase separation occurs (ds = 3 for classical 3D processes). (iii) The normalized first moment of the structure function, q1(t), exhibits a power time-dependence q1(t)/ t–a, the parameter a depending on the detailed mechanism of the aggregation of atoms.

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A. F. Craievich

(iv) The time dependence of the maximum of the structure function S(qm, t) is given 0 by S(qm, t) / ta with a0 = a.ds. All other moments and normalized moments of the structure function are also related by simple mathematical relations. A number of experimental investigations using small-angle (X-ray or neutron) scattering have demonstrated that the described dynamical scaling property also holds for phase separation processes occurring in many nanostructured materials, including glasses (Craievich and Sanchez 1981) and nanoporous xerogels (Santilli et al. 1995). Since the scattering intensity produced by very small primary particles (atoms) I1(q) is essentially constant within the small q range, all properties related to the time dependence of the structure function S(q) also apply to the time dependence of the experimental SAXS intensity function I(q).

Application (Example 1): Sintering of SnO2-Based Xerogels The theory described in the precedent section referring to phase separation processes in binary materials was applied to understand the structural evolution during isothermal treatment of nanoporous SnO2 xerogels studied by SAXS (Santilli et al. 1995). These nanoporous materials, after a short transient period, preserve their apparent density thus suggesting that the total fraction of porous volume remains constant during isothermal annealing. The series of SAXS curves displayed in Fig. 12a, corresponding to a SnO2-based xerogel isothermally annealed during increasing time periods at 400  C, exhibit a peak located at progressively decreasing q values. This feature is predicted by the statistical model described in the precedent section. The coincidence of all curves plotted as [S(q, t)q31/S2] vs. (q/q1) in Fig. 12b demonstrates that the dynamical scaling property (Eq. 38), theoretically derived for phase separation in simple binary systems, also applies to more complex processes such as the sintering of nanoporous xerogels.

Application (Example 2): Dynamical Scaling of Zirconia-Based Fractal Structures A demonstration of the dynamical scaling property for a system consisting of fractal zirconia-based aggregates embedded in a liquid matrix was reported by Lecomte et al. (2000). These authors analyzed the set of SAXS curves displayed in Fig. 10, which exhibit a maximum shifting progressively toward lower q for increasing periods of time. As pointed out before, the fractal dimension derived from the linear portions of the log I(q) vs. log q plots results D = 1.7. The same set of curves displayed in Fig. 10 was plotted in Fig. 13 using a [I(q/qm)  qds] vs. (q/qm) scale and setting ds = 1.7. In this analysis, the authors assumed that the first normalized moments q1 can be replaced as a reasonable approximation by the q-values associated to the maximum of the scattering

39

Small-Angle X-ray Scattering by Nanostructured Materials

a

1215

b [S(q).q13]/S2 (arb. units)

I(q) (arb. units) 8

1

4

0

0 0

0.05

0.10

q(Å−1)

0

2 q/q1

1

Fig. 12 (a) Scattering intensity curves corresponding to a nanoporous SnO2 based xerogel held at 400  C after increasing periods of time, from 4.5 min (bottom) up to 62 min (top). (b) Scaled structure functions [S(q, t)q31/S2] vs. q/q1 (Reprinted with permission from Santilli et al. (1995). Copyright 1995 by the American Physical Society)

Fig. 13 The same scattering intensity curves displayed in Fig. 10 replotted here as I(q)q1.7 m vs. q/qm, qm being the q value corresponding to the maximum of the scattering curves (Reprinted with permission from Lecomte et al. (2000). Copyright 2000 by the International Union of Crystallography)

I(q/qm).qm1.7 100 50 30 20 10 5 3 2 1

0.5

1

2

5 q/qm

10

20

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A. F. Craievich

curves qm. As it can be seen in Fig. 13 all scattering curves merge into a single scaled curve, this clearly demonstrating that the dynamical scaling property also applies to structural transformations of fractal aggregates. The results reported by Lecomte et al. (2000) referring to fractal structures demonstrated that the quotient of exponents a0 and a associated to the time dependences of the functions Sm(qm,t) and q1(t), respectively, would not be equal to the space dimension, ds = 3, but instead equal to the fractal dimension D. The described experimental results together with those mentioned in the preceding sections and others reported in the literature suggest that the statistical model derived for nanophase separation and particularly the dynamical scaling property of the structure function (Marro et al. (1975)) exhibit universal features, which provide a unified description of processes of structural coarsening in a wide variety of materials.

Grazing Incidence Small-Angle X-Ray Scattering Basic Concepts Thin films deposited on solid substrates such as those prepared by spin or dip coating and involving sol-gel transitions deserved the attention of many scientists because of their often interesting technological applications. These films usually have thicknesses ranging from about one nanometer up to a few microns. Since the structure of thin films supported by thick solid substrates cannot be studied by classical transmission SAXS, they are characterized by combining X-ray reflectivity and grazing incidence small-angle X-ray scattering (GISAXS). X-ray reflectivity measurements allow one to determine the thickness, average mass density, and surface roughness of thin films. Details of this experimental technique are not presented here. Readers interested on the basic concepts and applications of X-ray reflectivity are encouraged to consult the existing bibliography (for example Tolan 1999). Some thin films are heterogeneous at the nanometric scale. For example, thin films may be composed of a homogeneous matrix containing nanoclusters and/or nanopores, spatially correlated or not. Other materials consist of a homogeneous bulk volume with a thin layer close to their external surface containing buried nanoparticles. These nanoparticles are implanted by sputtering or plasma treatment or formed by nucleation and growth followed by atomic diffusion from supported thin films. GISAXS is usually applied to characterize nanostructured supported thin films and surface layers. Classical GISAXS experiments are performed using a flat sample, the incident beam hitting the sample surface at grazing incidence angles, αi, typically ~0.3 to 0.6 . The scattering patterns at small angles are recorded by a two-dimensional X-ray detector located at rather long distances from the sample, typically 1–3 m in synchrotron beam lines. Schematic views of the geometry of a GISAXS setup are shown in Fig. 14a, b. Notice that relevant angles in X-ray optics are measured with respect to the sample surface and not with respect to its normal as usual in classical optics.

39

Small-Angle X-ray Scattering by Nanostructured Materials

1217

Fig. 14 Schematic GISAXS setup. (a) Frontal view of the 2D X-ray detector taken along the direction of the incident beam. The distances that are measured in order to determine the three components of the scattering vector, qx, qy, and qz, associated to each detector pixel (ny, nz) (Eq. 39) are indicated. A narrow beam-stopper is usually vertically placed to avoid detector damaging by the strong reflected X-ray beam. (b) Lateral view indicating all relevant directions and angles. The GISAXS patterns can be recorded only above the sample horizon (dashed line) !

!

!

The three components of the scattering vector, q ¼ k  k 0 , associated to a scattered beam hitting a given detector pixel in GISAXS measurements (Fig. 14), are   qx ¼ ð2π=λÞ cos ψ cos αf  cos αi qy ¼ ð2π=λÞ sin ψ cos αf   qz ¼ ð2π=λÞ sin αf þ sin αi

(39)

The angles αf and ψ are determined from the vertical (dz) and horizontal (dy) distances (Fig. 14) as follows αf ¼ tg1½ðdz =LÞ  αi ψ ¼ tg1 d y =L

(40)

Notice that the detection plane of the 2D detector in real experiments is perpendicular to the incident X-ray beam and not parallel to the normal to sample surface, as

1218

A. F. Craievich

schematically shown in Fig. 14. Anyway, since the incidence angle αi is very small, the error in the angle αf associated to the use of this geometry can be neglected. For a X-ray beam propagating through a medium with refraction index n0 and hitting a flat interface with another medium with refraction index n, the angle of the refracted beam αr is determined by applying Snell law, ðn0 =nÞ ¼ ð cos αr = cos αi Þ. For an incident beam in vacuum ðno ¼ 1Þ or in a standard gas medium ðno  1Þ hitting a flat material surface, the refraction angle results: αr ¼ cos 1 ð cos αi =nÞ

(41)

For typical (small) incidence angles, αr is given as a good approximation by  1=2 αr ¼ α2i  2δ

(42)

where δ ¼ 1  n. Since the refraction index of any material for X-rays is slightly lower than 1(δ ~ 105), there is a critical value of the incidence angle, αc, for which the refracted beam propagates parallel to the sample flat surface. Substituting αr ¼ 0 in Eq. 42 the critical angle results: αc ¼ ð2δÞ1=2

(43)

Values of δ for any material composition and X-ray photon energy up to 30 KeV were reported by Henke et al. (1993). Let us now to describe the main features associated to specular reflection and refraction of an incident monochromatic X-ray beam hitting a flat sample surface under grazing incidence. For different incidence angles the following types of effects occur: (i) For αi < αc, the incident beam undergoes specular reflection at an exit angle αe = αi. (ii) For αi = αc, the refracted beam propagates parallel to the sample surface, i.e., αr = 0. (iii) For αi > αc, the refracted beam propagates inside the sample in a direction defined by the angle αr and amplitude t(αi) given by Snell law and Fresnel transmission function, respectively. The absorption of X-rays penetrating a flat sample produces a decrease in intensity of the incident beam described by the following basic equation: I ðdÞ ¼ I 0 eðμρd= sin αr Þ

(44)

where I0 is the intensity of the incident beam, μ is the mass absorption coefficient, ρ is the mass density, and d is the distance from the sample surface. The attenuation length (also named penetration depth) is the distance from the surface for which the intensity of the X-ray beam – penetrating into a given material

39

Small-Angle X-ray Scattering by Nanostructured Materials

1219

– becomes equal to I0/e. This distance is considered to be the approximate thickness of the layer probed in GISAXS measurements. Tables published by Henke et al. (1993) and a program accessible online in their web page yield the attenuation length as a function of the incidence angle, photon energy or wavelength, and sample chemical composition and mass density. If the incidence angle of the incoming X-ray beam is equal to or lower than the critical angle, an evanescent wave is formed whose penetration depth is only ~5 nm for typical materials. When a thicker layer is desired to be probed using an incident beam with same photon energy, the incidence angle αi should be set higher than αc. As an example, Fig. 15a displays the penetration depth of photons in three selected materials, SiO2, Al, and Ti, for a photon energy E = 8.04 KeV (corresponding to λCuKα = 1.542 Å) as a function of the incidence angle. Figure 15b shows the same functions for a photon energy E = 17.44 KeV (λMoKα = 0.707 Å). The curves displayed in Fig. 15a, b indicate that the Fig. 15 X-ray attenuation length or penetration depth as function of the incidence angle αi for three selected materials (SiO2, Al and Ti) for incident X-ray beams with two different wavelengths: (a) Cu λKα (1.542 Å) and (b) MoλKα (0.7071 Å). For an incidence angle below the critical angle the attenuation length is very short (a few nanometers)

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attenuation length below the critical angle αc (a few nanometers) is very small for the three selected materials, while for α > αc the attenuation length exhibits a monotonous increase for increasing incidence angles. Notice that the attenuation length is higher for less dense and lower Z materials. Thus, by adequate choices of incidence angle and photon energy, and depending on material composition and density, a wide range of thicknesses of nanostructured surface layers can be probed. An additional feature that is apparent in GISAXS patterns is named Yoneda peak (Yoneda 1963), which is associated to interference effects between the reflected and refracted waves. This peak appears at an exit angle αY ¼ αc with respect to the sample surface (Fig. 14a, b). The analysis of GISAXS results associated to nanostructured thin films and/or to surface layers is performed by fitting model functions to 2D experimental patterns. A reasonable model of a GISAXS function requires an initial guess of particle shape, size distribution, and structure function and should include the Fresnel transmission and reflection functions and the effects associated to Snell law. More detailed descriptions of GISAXS theory and applications were reported by Kutsch et al. (1997). Moreover, a recent review of modern applications of GISAXS, GISANS (grazing incidence neutron scattering) and grazing incidence X-ray and neutron wide angle scattering was published by Hexemer and MullerBuschbaum (2015).

Example of Application: Nanostructure of Thin Films Supported by Si Wafers A simple method for obtaining arrays of CoSi2 nanoplates endotaxially buried in a Si(001) single-crystalline wafer was reported by Kellermann et al. (2012). These authors demonstrated that thermally activated diffusion of Co atoms embedded in a Co-doped SiO2 thin film deposited on the (001) flat surface of a Si wafer promotes the formation of CoSi2 nanoplates buried inside the Si host. A transmission electron microscopy (TEM) study of this material indicated that the CoSi2 nanoplates exhibit a hexagonal lateral shape, are parallel to Si{111} crystallographic planes, have remarkably uniform sizes, and their lattices are coherently related to the host Si lattice. On the other hand, complementary analyses of TEM images showed the additional presence of a polydisperse set of spherical Co nanoparticles embedded in the supported SiO2 thin film. The model GISAXS function for a supported thin film containing an isotropic and dilute set of spherical nanoparticles with a radius distribution Nsph(R) is given by ð
     2   2 N sph ðRÞ:I 1 qx , qy , q~z , R dR Isph qy , qz / jtðαi Þj t αf

(45)

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where t(αi) and t(αf) are Fresnel transmission coefficients (Tolan 1999), I1,(qx, qy, qz, R) is given by Eq. 24 and refers to the SAXS intensity produced by spherical cobalt nanoparticles with radius R embedded in the silica thin film, and q~z is the z-component of the scattering vector considering that the incident beam scattered by nanoparticles is the refracted beam inside the sample. On the other hand, the GISAXS function associated to thin hexagonal CoSi2 nanoplates endotaxially buried in Si wafer, with their faces parallel to all Si{111} crystallographic planes, was modeled as (Kellermann et al. 2012; Kellermann et al. 2015): X   2     Ahex ðhklÞ αi , φ, qx , qy , q~z , L, T 2 I hex qy , qz / jtðαi Þj2 t αf  ðN hex =4Þ

(46)

hkl

where Nhex is the number of hexagonal nanoplates, Ahex(hkl) are the scattering amplitudes associated to regular hexagons oriented parallel to Si{111} crystallographic planes, with thickness T and lateral side L, φ is the azimuthal angle, and q~z is the component of the scattering vector in z direction inside the sample. The total GISAXS function associated to Co nanospheres embedded in the SiO2 thin film and CoSi2 nanohexagons buried in the Si wafer was modeled assuming independent contributions from both types of nano-objects, i.e.,       I total qy , q~z / ½I sph qy , q~z þ C:I hex qy , q~z 

(47)

where Isph and Ihex are given by Eqs. 45 and 46, respectively and C is an adjustable factor. In modeling the SAXS function, it was assumed that refraction effects are only produced at the interface between air and the SiO2 thin film. Because of the relatively low difference in density between the SiO2 thin film containing Co nanoparticles and the Si substrate, refraction effects associated to this interface were neglected. Kellermann et al. (2015) studied Co-doped SiO2 thin films deposited on silicon wafers with different surface orientations, namely, Si(001), Si(011), and Si(111), all of them previously heat treated at 750  C under identical conditions. Experimental 2D GISAXS patterns corresponding to different wafer orientations and the associated theoretical curves modeled by applying Eq. 47 are displayed in Fig. 16. The analysis of the experimental GISAXS results demonstrated that the sizes of the CoSi2 nanohexagons are functions of the crystallographic orientation of the Si substrate, the lateral size of the nanohexagons buried in Si(111) wafers being remarkably (~50%) larger than those grown inside the other two substrates, Si(011) and Si(001). The thickness of the platelets also varies for different Si substrate orientations from 2.8 nm for Si(001) up to 5.7 nm for Si(111). On the other hand, the spherical Co nanoparticles embedded in the SiO2 thin film exhibit average radii ranging from 0.6 nm for Si(011) up to 1.5 nm for Si(001).

Fig. 16 Experimental 2D GISAXS patterns corresponding to Co-doped SiO2 thin films deposited on (a) Si(001), (b) Si(111), (c) Si(011) flat wafers. Pictures (d–f) are the calculated patterns (Eq. 47) that best fit the patterns (a–c), respectively. The main axes of all elongated lobes in GISAXS patterns are perpendicular to Si{111} crystallographic planes (Reprinted from Kellermann et al., Phys. Chem. Chem. Phys. 2015; 17: 4945–4951)

1222 A. F. Craievich

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In conclusion, the GISAXS study reported by Kellermann et al. (2015) led to a complete low-resolution characterization of the nanostructures developed in Co-doped SiO2 thin films deposited on Si(001), Si(011), and Si(111) substrates.

Final Remarks A relevant issue omitted in this chapter is the experimental method that uses the properties of “anomalous” (or resonant) small-angle X-ray scattering (ASAXS), which is today widely applied thanks to the availability of tunable synchrotron X-ray sources (Goerigk et al. 2003). ASAXS is particularly useful for structural studies of biphasic materials with low contrast in electron density and also for analyses of complex multiphase systems. Additional information about SAXS theory is presented in the classical book authored by Guinier and Fournet (1955) and the book edited by Glatter and Kratky (1982). Another book dealing with SAXS, SANS, and light scattering was edited by Lindner and Zemb (1991). Instrumentation issues mainly focusing on SAXS using synchrotron radiation were described by Russell (1991), and SAXS/SANS studies of the structure and structural changes of biological macromolecules in solution were reviewed by Koch et al. (2003). A useful booklet for beginners was written by Schnablegger and Singh (2013). Besides the already mentioned GNOM software for SAXS data analysis (Svergun and Semenyuk 1991), new packages such as the recently developed SASFit (Kohlbrecher and Bressler 2014) are available to interested users. The amount of published articles based on the use of SAXS and SANS exhibited a fast increase during the past three decades (Craievich and Fischer 2010). This fast growth was primarily due to the increasing interest of scientists for studies of structural and physicochemical properties of nanomaterials. Other reasons that explain the observed strong growth in the annual number of published articles are: (i) the commercial availability of modern SAXS setups equipped with novel X-ray sources, focusing and collimating optics and fast high-resolution 2D detectors, (ii) the development of new theoretical approaches and numerical methods for data analysis, (iii) the increasing availability of powerful computers, and (iv) the opening of new small-angle scattering beam lines in many synchrotron and neutron laboratories around the world.

Appendix: Experimental Issues Basic Comments Monochromatic X-ray beams are characterized by their photon energy E or wavelength λ, both related by λ = hc/Ε, where h is the Plank constant and c is the speed of light in vacuum, i.e., λ(Å) = 12.398/E(KeV). The wavelengths of typical monochromatic beams

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used in SAXS experiments are within the range 0.6–2.0 Å circa (i.e., photon energies ranging from ~6 to ~20 KeV). The X-ray beams produced by synchrotrons or typical commercial sources are usually monochromatized by quartz, germanium, or silicon single crystals, which yield incident beams with very narrow pass-bands (Δλ/λ < 103). Considering, for example, a typical SAXS experiment with an X-ray wavelength λ = 1.542 Å (λCuKα), a sample-to-detector distance D = 1 m, a beamstopper with a diameter ϕ1 = 5 mm and a circular 2D detector with a diameter ϕ2 = 150 mm, and remembering that q ¼ ð4π=λÞ sin θ  ð2π=λÞ:2θ for low q, the range of scattering angles to be covered results 0.14 < 2θ < 4.3 , and the corresponding minimum and maximum q values are 0.01 Å1 and 0.30 Å1, respectively. Different lower and upper q limits can be reached by selecting adequate beam collimation, sample-to-detector distances and/or X-ray wavelengths. The choice of the experimental q range depends on the sizes of the nanoparticles to be studied. X-ray beams for SAXS experiments are produced by classical sealed X-ray tubes, rotating anode X-ray generators and synchrotron sources. Synchrotron radiation sources are often preferred because they provide powerful, continuously tunable and well collimated X-ray beams. Another closely related experimental technique often used for same or similar purposes is small-angle neutron scattering (SANS), its basic theory being essentially the same as that developed for the SAXS technique.

Choice of Sample Thickness Classical SAXS experiments are performed in transmission mode and usually under normal incidence. The first step for planning SAXS experiments is to determine the sample thickness that maximizes the scattering intensity for a given material and photon energy. The SAXS intensity produced by any material with arbitrary structure, as a function of sample thickness t, is given by I ðtÞ / teμρt

(48)

where ρ is the mass density and μ the mass X-ray absorption coefficient, which is a function of chemical composition of the material and photon energy. The absorption coefficient can be obtained from tables published by Henke et al. (1993) or by using an online program accessible in their web page. Examples of the function defined by Eq. 48 are plotted in Fig. 17 for three different materials. The optimum thickness tmax corresponding to the maximum of the I(t) function is tmax ¼ ðρμÞ1

(49)

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Fig. 17 (a) Examples of scattering intensities in arbitrary units as functions of sample thickness for an incident X-ray beam with a wavelength λCuKα = 1.542 Å, corresponding to different selected materials: Cu, SiO2, and H2O, whose optimum thicknesses tmax are 22 μm, 132 μm, and 1.00 mm, respectively

This implies that the transmittance of samples with optimum thickness is T ¼ ðI transmitted =I incident Þ ¼ e1 ¼ 0:37:

(50)

Notice that the tmax values determined by Eq. 49 are just a guide for a convenient choice of sample thickness. However, it is always advisable to avoid the use of very thick or very thin samples which would lead to high absorption and low probed volumes, respectively, both yielding weak scattering intensities. For samples containing large fractions of high Z atoms the optimum thicknesses could be extremely low using CuλKα photons (E = 8.04 KeV). For these materials X-ray beams with higher photon energy should be employed. On the other hand, in order to minimize fluorescence effects, the use of beams with photon energies above and close to absorption edges of sample elements should be avoided.

Subtraction of Parasitic Scattering Before further analysis of experimental SAXS results, a pretreatment of rough data is ! required. For anisotropic 2D SAXS patterns, the vector q associated to each detector pixel is calculated. For isotropic 2D SAXS patterns, the scattering intensity is defined as a function of the modulus of the scattering vector, which is determined by circular averaging.

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In order to subtract the parasitic scattering intensity produced by slits, cell windows, and air, two SAXS patterns should be recorded: (i) the total scattering
 ! intensity (from sample plus parasitic scattering) defined by the counting rate RT q
 ! and (ii) the parasitic scattering intensity given by the counting rate RP q recorded under same experimental conditions but without sample. The scattering intensity exclusively related to the sample is given by
 h

  i h

  i ! ! ! R q ¼ RT q  RD =T SþW  RP q  RD =T W

(51)

where RD is the counting rate associated to the detector noise, TS+W is the transmittance of the sample and cell thin windows, and TW is the transmittance of the empty sample cell. For solid samples placed in a windowless holder, we have TW = 1. Often the counting rate associated to parasitic scattering for macromolecules in dilute solution is determined with the sample cell filled with same buffer, thus under this condition the scattering intensity due to statistical density fluctuations in the solvent is also subtracted. When SAXS experiments are conducted using synchrotron beam lines with continuously decreasing electronic current, the effects of time variation of the intensity of the incident X-ray beam should be properly accounted for.

Correction of Smearing Effects The use of X-ray incident beam with rather large cross-section and/or X-ray detectors with large pixel size may produce serious smearing effects on the SAXS curves. However, most of the modern commercial setups and synchrotron beam lines provide an incident beam with pinhole-like cross-section and use X-ray detectors with very small pixel size, thus often making mathematical desmearing procedures unnecessary. When using commercial setups yielding an incident X-ray beam with large cross section (for example a beam with linear cross-section), two approaches can be applied for quantitative analyses: (i) fitting the theoretical model of SAXS curve to previously dismeared experimental functions or (ii) fitting the previously smeared theoretical model of SAXS curve to the experimental function. Since mathematical desmearing of experimental SAXS patterns leads to results with rather high statistical noise, the second procedure is generally preferred.

Determinations of SAXS Intensity in Relative and Absolute Units For pin-hole collimation of the incident beam, the counting rate


 ! R q

corresponding to the X-ray photons scattered by the sample is proportional to the

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 ! I q function used along this chapter. Thus Eq. 51 directly yields the scattering intensity in relative scale or arbitrary units, to which model functions are fitted after adequate scaling. However, the SAXS intensity given in absolute scale provides additional information that is often useful for detailed structural characterization. The typical scattering intensity function in absolute scale is the differential scattering cross-section per unit volume (dΣ/dΩ). This function is
related
  to
theSAXS intensity !

!

I q , which was used along this chapter, by ðdΣ=dΩÞ q

!

¼ I q :r 2e =V.

For SAXS measurements using pin-hole collimation (i.e., with a point-like incident beam cross-section), the differential scattering cross-section per unit volume is given by
 ! dΣ
! R q =η q ¼ dΩ I 0 V:ΔΩ

(52)


 ! where R q is the photon counting rate, η is the detector efficiency, I0 is the photon flux of the incident X-ray beam (number of photons per unit cross-section.second), V is the probed sample volume, and ΔΩ is the solid angle associated to the surface area of the detector pixel. The usual unit for the differential scattering cross-section per unit volume is cm1. Equation 52 can also be written as
 !
 R q :L2 dΣ ! q ¼ dΩ R0 ts Δa

(53)

where R0 is the counting rate (number of photons/second) corresponding to the total incident beam, ts is the sample thickness, Δa is the surface area of the detector pixel, and L is the sample-to-detector It is assumed in Eq. 53 that the efficiency of
distance.  !

the detectors that records R q and Ro are identical. When different detectors are
 ! used for the measurements R q and Ro, the counting rates should be properly normalized to equivalent efficiencies. Equation 53 is usually applied to plate-shaped solid samples or to liquids contained in cells with parallel thin windows for entrance of the incident X-ray beam and exit of the scattered photons. Determinations of SAXS intensity in absolute units associated to powdered samples or liquid samples contained in cylindrical capillaries are also possible but their evaluation is less precise (Fan et al. 2010). Since the measurement of Ro is in practice difficult using standard detectors, the differential scattering cross-section per unit volume of solid materials is generally

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determined by means of an independently calibrated sample, such as Lupolen or glassy carbon (Fan et al. 2010). In order to determine the differential scattering cross-section per unit volume associated to colloidal particles embedded in a liquid medium, it is also recorded – under the same experimental conditions – the SAXS intensity produced by statistical density fluctuations in water. The differential scattering cross-section per unit volume of water ðdΣ=dΩÞH2 O – which is a isotropic and constant function at small q – is given by (Guinier and Fournet 1955):

dΣ dΩ

ðq ! 0Þ ¼ ðN H2 O ne r e Þ2 kTβ

(54)

H2 O

where N H2 O is the number of water molecules per unit volume, ne is the number of electrons per water molecule, k is the Boltzmann constant, T is the absolute temperature, and β is the isothermal compressibility of water at room temperature. Since all parameters in Eq. 54 are known, the differential scattering cross-section per unit volume of water can be written as

dΣ dΩ



ðq ! 0Þ ¼ 1:65:102 cm1

(55)

H2 O

If the counting rate associated to a isotropic liquid sample (for example proteins in liquid buffer), [R(q)]sample, and that corresponding to water, RðqÞH2 O, are determined under same experimental conditions, the differential scattering cross section per unit volume of the studied sample is given by ½RðqÞsample dΣ ðqÞ ¼ 1:65:102 cm1 dΩ < Rð q Þ H 2 O >

(56)

where < RðqÞH2 O > is an average value taken within the small q range over which the counting rate is approximately constant. If SAXS measurements corresponding to sample and water are conducted under different experimental conditions, adequate corrections should be applied. Additional details on this matter were reported by Fan et al. (2010). Acknowledgments The author thanks the staff of the National Synchrotron Radiation Laboratory (LNLS), Campinas, Brazil, where the experimental parts of most of the SAXS investigations reported in this chapter were conducted; G. Kellermann and C. Huck-Iriart for their useful remarks; and H. Fischer for his help for figures preparation.

References Beaucage G, Ulibarri T, Black EP, Shaeffer DW. Chapter 9. Multiple size scale structures in silicasiloxane composites studied by small-angle scattering. In: Mark JE, Lee CYC, Bianconi PA, editors. Hybrids organic–inorganic composites, vol. ACS series 585. Washington, DC: American Chemical Society; 1995. p. 97–111.

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Cahn JW. Phase separation by spinodal decomposition in isotropic systems. J Chem Phys. 1965;42:93. Ciccariello S. The leading asymptotic term of the small-angle intensities scattered by some idealized systems. J Appl Crystallogr. 1991;24:509–15. Ciccariello S, Schneider JM, Schonfeld B, Kostorz G. Illustration of the anisotropic Porod law. J Appl Crystallogr. 2002;35:304–13. Craievich AF, Alves OL, Barbosa LC. Formation and growth of semiconductor PbTe nanocrystals in a borosilicate glass matrix. J Appl Crystallogr. 1997;30:623–7. Craievich AF, Fischer H. Quantitative analysis and relevant features of the scientific literature related to SAXS and SANS. J Phys Conf Ser. 2010;247:012003. Craievich AF, Sanchez JM. Dynamical scaling in the glass system B2O3–PbO–Al2O3. Phys Rev Lett. 1981;47:1301311. Dahmouche K, Santilli CV, Pulcinelli SH, Craievich AF. Small-angle X-ray scattering study of sol–gel-derived siloxane-PEG and siloxane-PPG hybrid materials. J Phys Chem B. 1999;103:4937–42. Debye P, Bueche AM. Scattering by an inhomogeneous solid. J Appl Phys. 1949;20:51525. Fan L, Degen M, Bendle S, Grupido N, Ilavsky J. The absolute calibration of a small-angle scattering instrument with a laboratory X-ray source. J Phys Conf Ser. 2010;247:012005. Glatter O, Kratky O, editors. Small-angle X-ray scattering. London: Academic; 1982. Goerigk G, Haubold HG, Lyon O, Simon JP. Anomalous small-angle X-ray scattering in materials science. J Appl Crystallogr. 2003;36:425. Guinier A, Fournet G. Small-angle scattering of X-rays. New York: Wiley; 1955. Henke BL, Gullikson EM, Davis JC. X-ray interactions: photoabsorption, scattering, transmission and reflection at E = 50-30000 eV and Z = 1-92. Atomic Data and Nuclear Data Tables. 1993;54:181–342. http://henke.lbl.gov/optical_constants/ Hexemer A, Muller-Buschbaum P. Advanced grazing incidence techniques for modern soft-matter materials analysis. IUCrJ. 2015;2:106–25. Kellermann G, Montoro LA, Giovanetti LJ, Santos Claro PC, Zhang L, Ramirez AJ, Requejo FG, Craievich AF. Formation of an extended CoSi2 thin nanohexagons array coherently buried in silicon single crystal. Appl Phys Lett. 2012;100:063116. Kellermann G, Montoro LA, Giovanetti LJ, dos Santos Claro PC, Zhang L, Ramirez AJ, Requejo FG, Craievich AF. Controlled growth of extended arrays of CoSi2 hexagonal nanoplatelets buried in buried in Si(001), Si(011) and Si(111) wafers. Phys Chem Chem Phys. 2015;17: 4945–51. Koch MHJ, Vachette P, Svergun DI. Small-angle scattering: a view on the properties, structures and structural changes of biological macromolecules in solution. Q Rev Biophys. 2003;36:147–227. Kohlbrecher J, Bressler I. Software package SASfit for fitting small-angle scattering curves. 2014. http://kurweb.psi.ch/sans1/SANSSoft/sasfit.html Kustch B, Lyon O, Schmitt M, Mennig M, Schmidt H. Small-angle X-ray scattering experiments in grazing incidence on sol–gel coatings containing nano-scaled gold colloids: A new technique for investigating thin coatings and films. J Appl Crystallogr. 1997;30:94956. Lebowitz JL, Marro J, Kalos MK. Dynamical scaling of structure-function in quenched binaryalloys. Acta Metall. 1982;30:297–310. Lecomte A, Dauger A, Lenormand P. Dynamical scaling property of colloidal aggregation in a zirconia-based precursor sol during gelation. J Appl Crystallogr. 2000;33:496–9. Lindner P, Zemb T, editors. Neutron, X-ray and light scattering. Amsterdam: North Holland; 1991. Marro J, Boltz AB, Kalos MH, Lebowitz JL. Time evolution of a quenched binary alloy. II. Computer simulation of a three-dimensional model system. Phys Rev B. 1975;12:2000–11. Meakin P. In: Stanley HE, Ostrowsky N, editors. On growth and form. Boston: Martinus Nijhoff; 1986. p. 111–35. Porod G. Chapter 2: General theory. In: Glatter O, Kratky O, editors. Small-angle X-ray scattering. London: Academic; 1982. Riello P, Minesso A, Craievich AF, Benedetti A. Synchrotron SAXS study of the mechanisms of aggregation of sulfate zirconia sols. J Phys Chem B. 2003;107:3390–9.

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Ruland W. Small-angle scattering of 2-phase systems. Determination and significance of systematic deviations from Porod’s law. J Appl Crystallogr. 1971;4:70. Russell TP. Chapter 11: Small-angle scattering in synchrotron radiation sources. In: Brown GS, Moncton DE, editors. Handbook on synchrotron radiation, vol. 3. Amsterdam: North Holland; 1991. Santilli CV, Pulcinelli SH, Craievich AF. Porosity evolution in SnO2 xerogels during sintering under isothermal conditions. Phys Rev B. 1995;51:8801–9. Schnablegger H, Singh Y. The SAXS guide. Graz: Anton Paar GmbH; 2013. Shull CG, Roess LC. X-ray scattering at small angles by finely-divided solids. I. General approximate theory and applications. J Appl Phys. 1947;18:295–307. Silva NJO, Dahmouche K, Santilli CV, Amaral VS, Carlos LD, V. BZ, Craievich AF. Structure of magnetic poly(oxyethylene)-siloxane nanohybrids doped with Fe-II and Fe-III. J Appl Crystallogr. 2003;36:961–6. Svergun DI. Determination of the regularization parameter in indirect-transform methods using perceptual criteria. J Appl Crystallogr. 1992;25:495–503. Svergun DI. Restoring low resolution structure of biological macromolecules from solution scattering using simulated annealing. Biophys J. 1999;77:2879–86. Svergun DI, Semenyuk A. Small-angle scattering data processing using the regularization technique.1991. www.embl-hamburg.de/biosaxs/gnom.html Teixeira J. Small-angle scattering by fractal systems. J Appl Crystallogr. 1988;21:781–5. Tokumoto MS, Pulcinelli SH, Santilli CV, Craievich AF. SAXS study of the kinetics of formation of ZnO colloidal suspensions. J Non-Cryst Solids. 1999;247:176–82. Tolan M. X-ray scattering from soft-matter thin films. Berlin: Springer; 1999. Yoneda Y. Anomalous surface reflection of X-rays. Phys Rev. 1963;131:2010–3.

X-ray Absorption Spectroscopy Studies on Materials Obtained by the Sol-Gel Route

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Francesco d’Acapito

Contents Introduction to X-ray Absorption Spectroscopy (XAS) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Theory . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Data Collection and Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Data Analysis Procedures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Applications of XAS to Sol-Gel Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Materials for Optics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Materials for Catalysts and Sensors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Electrodes for Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Perspectives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1232 1232 1237 1239 1241 1241 1244 1247 1249 1251 1252

Abstract

This chapter reports on the use of X-ray Absorption Spectroscopy in the characterization of materials obtained by the Sol-Gel method. Firstly, an introduction to the theoretical bases of the technique is given followed by a brief description of the most relevant data collection schemes and data analysis. Successively, a collection of recent relevant experiments is presented putting in evidence, case by case, the peculiar information retrieved thanks to this technique. These studies are classed following the kind of material or its use: materials for optics, nanoparticles, materials for catalysis and sensors, and materials for batteries. Eventually, a brief overview is given on the future perspectives of the technique in terms of new data collection methods and data analysis. A rich list of bibliographic references completes this contribution.

F. d’Acapito (*) CNR-IOM-OGG c/o European Synchrotron Radiation Facility, LISA CRG, Grenoble, France e-mail: [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_31

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Introduction to X-ray Absorption Spectroscopy (XAS) Theory

Absorption Coefficient µ (Arb. Units)

X-ray Absorption spectroscopy (XAS) (Lee et al. 1981) is an experimental technique that permits the quantitative determination of structural parameters around chosen atomic species and to obtain a description of the electronic structure. In particular, it derives this information from the small oscillations that appear in the absorption coefficient as a function of energy just above the edge step from a deep core state (1s or 2s, 2p). These oscillations appear only when the absorbing atom is regularly surrounded by other atoms as shown in Fig. 1. XAS permits the determination of the number of neighbors N and distance R from the atom that absorbs the X-ray photon (hereafter defined as the absorber) with an accuracy of about 10% for N and 1% for R. Moreover, it is chemical selective as the absorption edges of elements are well spaced in energy from each other (for edge energy values refer to data in Bearden and Burr (1967)). The presence of small oscillations above the absorption edge was reported much earlier but only in the early 1970s (Sayers et al. 1971), and with the successive advent of synchrotron radiation facilities (Kincaid and Eisenberger 1975), XAS theory has been fully understood and developed (Rehr and Albers 2000), and the technique has been used in a variety of fields like solid state physics, materials science, environmental science, chemistry, structural biology, and archaeometry as can be seen in a series of publications on this technique (Köningsberger 1988; Bunker 2010). The physical origin of the μ oscillations above the edge is the modification, due to the neighboring atoms, of the final state of the electron emitted by the atom that

10800 11000 11200 11400 11600 11800 12000 12200 12400 Energy (ev)

Fig. 1 Example of a XAS spectrum: the Ge-K edge of crystalline Ge (Data collected at the GILDA beamline, European Synchrotron Radiation Facility)

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absorbs the incoming X-ray photon (hereafter called “Central” or “Absorber” atom) (Lee et al. 1981; Stern 1988). The X-ray absorption cross-section, Σ, of an atom !

interacting with a photon described by a dipole transition operator D is given by (Natoli and Benfatto 1986): Σ ¼ 4π2 Eα

X D  ! ! E2  i p  D f  δðE  Ef þ E0 Þ

(1)

f

where E is the photon energy, α the fine structure constant, EF is the photoelectron energy, and E0 is the threshold   energy. hii and hfi are, respectively, the initial and !

final electron states, and p is the electron momentum operator. The initial state hii

is usually an atomic core state (K or L) so it is strongly localized around the central atom site: This means that the space integrals in Eq. 1 are calculated in a restricted region around the absorber. In the case of the isolated atom, h f | is an outgoing qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 0Þ spherical wave where k is the photoelectron wavevector k ¼ 2m ðEE ℏ . In this case, hfi is a smooth and structureless function of E, and a similar behavior will be found in the absorption cross section Σ0. When a neighbor (Backscattering atom) is placed near to the central atom, a part of the outcoming electron wave will be backscattered by the neighbor: hf i will be the sum of the original outgoing wave plus a modified part hΔf i, coming from the waves backscattered by the neighbors. The backscattered wave will have an amplitude A(k, Rj) and a phase φ(k, Rj) which depend on the details of the scattering potential. The interference between these two terms in the absorber site gives rise to the oscillations observed in the experimental data. The oscillatory part χ of the signal is defined considering Σ0 (containing hfi) and 0Þ Σ (containing hf þ Δf i) and taking χ ¼ ðΣΣ Σ0 . The explicit expression for χ in the case of an “S” state excitation and considering only single scattering events is (Lee et al. 1981): χ¼

X j

χj ¼

X j

3S20

2 2       Nj  ^ j 2 (2) A k, Rj sin 2kRj þ φ k, Rj þ 2δc ðkÞ ek σj e2Rj λ ^e  R 2 kRj

Here the sum is extended over the various coordination shells indexed by j, each containing Nj identical atoms at a distance Rj from the central one. The backscattering amplitude and phase, A and φ, are expressed in terms of the bond length Rj ! (curved wave approximation). In Eq. 2, e is the field polarization unit vector and the related term expresses the dependence of χ on the beam polarization. This depen   ! ! dence comes from the p  D operator and is of particular interest for single crystal samples, where the absorption coefficient depends on the sample orientation. This effect is different if {K, LI} or {LII, LIII} edges are considered. In the former case, the relation between the amplitude N of a single bond making an angle δ with the polarization vector is NK, LI ¼ 3  cos2 δ, whereas in the latter case, it is

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¼ 0:7 þ 0:9  cos2 δ (Iwasawa 1997). Conversely, in the case of isotropic   ! ! 2 samples (polycrystalline powders, amorphous, liquids,. . .), the e  R j term is NLII ,

LIII

replaced by its angular average, 1/3. In the χ(k) formula shown above, other terms are added to the simple scattering approximation in order to account for two damping processes. One is the limited lifetime of both the photoelectron and the core-hole. Indeed, the excited atom keeps this state for a limited time, after which an electron from the upper state fills the hole in the core state, destroying the extended X-ray absorption fine structure (EXAFS) signal. This problem is particularly severe when working on edges of heavy elements at high (E > 40 keV) energies (Borowski et al. 1999) as the broadening due to the core-hole limited lifetime can be several eV (Krause and Oliver 1979). Also, when the emitted photoelectron undergoes momentum transfer or inelastic scattering with other electrons, the EXAFS signal is cancelled. All this can be accounted for by adding an imaginary part to the scattering potential (Chou et al. 1987) and is contained in the exp(rj/λ) term. As λ is of the order of a few Å, EXAFS signals coming from long paths are rapidly damped and this makes the technique sensitive only to the local atomic structure around the absorber. Simultaneously with the electron emission, other electrons can also be excited in the central atom, both to bound states or continuum states (Lee and Beni 1977), as shown by photoemission experiments. This reduction of the overall amplitude is accounted for by the S02 term. The physical process at the origin of the EXAFS signal is interference, so it strongly depends on the spatial disorder of the back-scattering centers. The disorder can be originated by thermal vibrations or by random positions of the neighbor atoms in amorphous systems. Considering this, the χj function should be calculated as the integral of χj(R), weighted by the Pair Distribution Function (PDF) relative to the i-th shell (Benfatto et al. 1989). If we suppose a small disorder and a gaussian pair distribution function with mean square displacement σ2j, the result of the integral is the exp(k2σj2) term in Eq. 2 (Beni and Platzmann 1976). Note that this term, although similar to the Debye-Waller factor in X-ray diffraction, has here a different meaning. Whereas in the latter case, σ is related to the variation of the atomic position with respect to the ideal lattice site, in EXAFS it means the relative displacement of the neighbor relative to the absorber position. If thermal disorder dominates, the values of σj can be calculated from the dynamic properties of the lattice under study, as shown in Beni and Platzmann (1976) and Sevillano et al. (1979). When the atomic oscillations become large, and this is the case, namely, for lattices at temperatures near or above their Debye temperature or in case of strongly disordered systems, the pair distribution functions also acquires a nonsymmetric character with respect to the distribution maximum, i.e., they become skewed at higher r values. This is due to the atomic interaction potential, which generally exhibits a Lennard-Jones-like behavior. The asymmetry of the PDF leads to unphysical results if data are treated in the gaussian approximation, typically a contraction of the distances and a drop in the coordination numbers, as shown in Eisenberger and Brown (1979). In this case, an analysis based on a more detailed

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model for the PDF is needed, by taking into account higher cumulants of the distribution (Tranquada and Ingalls 1983) or introducing ad hoc PDFs (Filipponi et al. 1995; Filipponi and DiCicco 1995; Kuzmin 1997) for the construction of the χ function. So far we have considered the photoelectron interactions in the solid as a series of single scattering events with the surrounding atoms. In a more general view, we should take into account processes where the photoexcited electron is scattered by several neighbors before returning to the central atom. While the scattering amplitude dies with the path length, the number of possible paths increases and it comes out that multiple scattering (MS) phenomena are not negligible in the interpretation of an X-ray Absorption Spectrum (Lee and Beni 1977) and extend up to several tens of eV above the edge (Benfatto et al. 1986). The calculation of the photoelectron final state hf i, accounting for multiple scattering events, can be carried out in several ways, like in the scattering approach, the band-structure approach, or the Greenfunction approach (Natoli and Benfatto 1986), all leading exactly to the same result. The X-ray cross section is expressed in matrix form by Natoli and Benfatto (1986): "

# i0, 0 X h 1 1 σ ¼ σ0 ℑ ðI  TGÞ T  lm, lm ð2L0 þ 1Þ sin2 δ0l m " # X X 1 0, 0 ℑ ½ðTGÞn Tlm, lm σ0 ℑ 2 0 ð2L0 þ 1Þ sin δl m n

(3)

Here, G is the so-called Propagator matrix and contains all the geometrical details of the cluster surrounding the absorber, whereas T is the “Scattering matrix” and it contains the details of the scattering potentials. L0 is the angular moment of the initial state, and σat is the (structureless) atomic cross section. The (1 + GT)1 term can, in some cases, be replaced by its series expansion, giving rise to the so-called multiple scattering series approximation of σ (right side of Eq. 3). The terms with n = 2 in the expansion correspond to single scattering events (photoelectron scattered by the neighbor and back to the absorber), whereas terms with n = 3, n = 4, . . . correspond to events where the photoelectron is scattered 2, 3, . . . times by other neighbors, before reaching back the absorber. Examples of single and multiple scattering paths are given in Fig. 4. The multiple scattering expansion is valid provided that the related series is convergent. As the G matrix elements depend on 1/k or, equivalently, pffiffiffiffiffiffiffiffiffiffiffiffiffiffi on 1= E  E0 (whereas the elements of the T matrix depend on the chemical nature of the scatterer), the convergence criterium is usually satisfied for photoelectron energies above a few tens of eV and depends on the system under investigation (Natoli and Benfatto 1986). The standard EXAFS formula shown in the first paragraph can be obtained from the MS expansion, taking L0 = 0 and cutting the series at the n = 2 term. For the higher terms, it can be demonstrated (Zabinsky et al. 1995) that, for any given path i, they can be reduced to expressions similar to that shown in Eq. 2 with suitable amplitudes Ain(k, Ri) and phases φin(k, Ri), where now Ri is half the total path length. In the example of Fig. 4, the path in (c) is associated to a given χi and the path

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in (d) to a different χj. By summing all these partial χi, the total signal can be reproduced. Debye Waller factors related to Multiple Scattering paths have been defined via a generalized formula as presented in Poiarkova and Rehr (1999). Using a different approach (Filipponi and DiCicco 1995), the contribution to the total χ of a given atomic arrangement, called γi, can also be parametrized similarly to Eq. 2. When approaching the absorption edge, the photoelectron mean free path becomes appreciably large (as well as its wavelength) and a larger number of paths contribute to the total signal. There exists a limit beyond which the MS expansion no longer converges, so it is not possible to calculate the cross section in this way. The 1/(1  GT) matrix has to be calculated explicitly and this is the so-called Full Multiple Scattering method (Tyson et al. 1992). This region is called XANES (X-ray Absorption Near Edge Structure) and it contains in principle threedimensional information on the absorber site (Fig. 2). Apart from the complex spectrum simulations, XANES can be easily used for the determination of the valence state and local symmetry of the element under investigation prior to comparison with known compounds. The valence state is derived from the energy position of the edge (Cramer et al. 1976), the more oxidized states corresponding to higher edge energy values. The local symmetry can be derived from the intensity of the small peak appearing before the edge (that in the case of K

Fig. 2 Partition of a typical XAS spectrum in the EXAFS and XANES regions, here for the spectrum of hematite (Fe2O3) at the Fe-K edge. In the EXAFS region (roughly above 50 eV from the edge), only Single Scattering (SS) or a mixture of Single and a limited number of Multiple Scattering paths contribute to the spectrum. In the XANES region, all the possible paths contribute to the spectrum and the Full Multiple Scattering (FMS) approach is needed to reproduce the spectrum

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edges of the 3d metals are due to partially forbidden 1s-3d transitions) with tetrahedral environments exhibiting more intense peaks (Galoisy et al. 2001).

Data Collection and Analysis The absorption coefficient, μ, of a sample as a function of the energy is defined as:

 Φ 1 ð EÞ μðEÞ ¼ ln Φ 0 ð EÞ

(4)

where Φ0 and Φ1 are the photon fluxes before and after the sample, respectively (Fig. 3a). The fluxes are normally measured by an ionization chamber, usually filled with low-pressure gas. There are different methods to measure the absorption coefficient, depending on the nature of the sample being investigated. As a thumbrule, when the atomic species of interest contributes significantly to the total absorption of the sample, it is convenient measuring the intensity transmitted after the sample (transmission mode). This is commonly the case of systems with a high number of absorbing centers, like bulk, heavy (high atomic number Z) elements. The measure of the transmitted intensity is carried out by a second ionization chamber (Lee et al. 1981).

Fig. 3 Experimental setup for the direct (a) and indirect (b) measurement modes of the X-ray absorption coefficient of a sample

a

Ic 1

Sample

Ic 0

Transmission Mode

b

X-ray or Electron Detector

Ic 0

Sample Fluorescence (or Electron detection) Mode

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F. d’Acapito

In some cases, e.g., samples with a low concentration of absorbers, it is better to measure indirectly μexp, looking at the processes usually coupled to the X-ray absorption. Indeed, the photoionization process creates a hole in the core state that is suddenly filled by another electron. This occurs through a couple of competitive processes: a radiative one, where an electron “falls” from a higher energy state to the ionized one emitting a photon (fluorescence). Alternatively, hole filling can be accompanied by the emission of fast electrons to balance the energy (Auger processes). The higher the Z of the involved atom, the more probable is the fluorescence process, with equal probability occurring at roughly Z = 30. It must be noted that the emission of the photon occurs at a well-determined energy, depending on the atom and on the level involved. It is thus easy to separate the desired signal from the background. Exploiting the first effect described above is called the fluorescence collection mode (Jaklevich et al. 1977) (Fig. 3b). In this case, an energy-selective detector (namely, solid state detectors such Li:Si, Si Drift, High Purity Ge, . . .) is used to separate the fluorescence from the background (coherent and incoherent scattering), or fluorescence from other elements. This method is particularly well adapted for the analysis of diluted samples but leads to spectrum distortions when applied to concentrated samples. In this case either data-correction routines have to be applied (Troger et al. 1992) or data have to be collected at grazing exit angle (Pfalzer et al. 1999). Recently, a novel detection scheme for fluorescence XAS on concentrated samples has been proposed (Achkar et al. 2011) where the absorption signal of an atom A is collected by recording the modulations of the fluorescence yield of a lighter atom B while scanning the energy through the A absorption edge. A method that gained considerable popularity in the latest years is the High Energy Resolution Fluorescence Detection (HERFD). Indeed, as shown in H€am€al€ainen et al. (1991), collecting the fluorescence yield from a sample with an energy resolution lower than the core-hole linewidth produces XANES spectra no longer broadened by the initial state, so permitting to evidence structures in the spectrum barely visible with the conventional data collection. This technique is called High Energy Resolution Fluorescence Detection (HERFD) and results to be particularly effective in the analysis of the L edges of 5d metals where the typical broadening of the states is of the order of 5–10 eV. The detector is usually a multiple crystal analyzer with a limited solid angle acceptance, so this technique can be carried out only on very intense sources. A description of a complete spectrometer can be found in Glatzel et al. (2005) and Rovezzi and Glatzel (2014). The drawback is that using crystal analyzers greatly reduces the effective collection solid angle for the fluorescence so the total efficiency (flux on the detector) results to be greatly reduced with respect to the conventional method with a considerable increase of the minimum concentration limits. Alternatively to fluorescence, the total electron yield from the sample due to cascades initiated by the Auger processes can be detected (Citrin et al. 1978). The signal is measured with a channeltron detector or simply by measuring the photoionization current from the sample. This method has the peculiarity of probing a few thousand Å under the sample surface (due to the limited electron escape depth) and

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can be useful in studying macroscopically layered structures, e.g., ion-implanted materials. The other way to study surfaces is based on the x-ray total reflection: When the x-ray beam impinges with a sufficiently small angle on the sample, the x-ray penetration is limited to a few Å (Parratt 1954). In this case, the signal can be collected both from the fluorescence yield from the sample or from the reflected beam, using detectors described above such as photodiodes or scintillation detectors (Heald et al. 1988). In this setup, the ultra-high vacuum environment is not needed and it can be exploited to investigate liquid surfaces buried interfaces or gas-solid interfaces. An example of a typical experimental apparatus for total reflection EXAFS can be found in d’Acapito et al. (2003 and references therein) and consists of an accurate sample positioning stage with detectors for the impinging and reflected beam as well as for the fluorescence from the sample. The main limitation of this procedure is the need of sufficiently long (a few cm) and smooth (microscopeslide grade) samples. This method, enhancing the signal from the surface thin film with respect to the substrate, is particularly indicated in the study of thin layers like those obtained by sol-gel (d’Acapito et al. 2008). By exploiting the luminescent properties of the samples, it is possible to derive interesting results using the X-ray excited optical luminescence (XEOL) technique. XEOL is a particular data collection technique (Rogalev and Goulon 2002; Sham and Rosenberg 2007) which derives the absorption coefficient μ from the optical luminescence that follows the absorption of an X-ray photon in particular materials. The electrons emitted in each absorption event produce a shower of lower energy electrons (that extends far from the X-ray absorbing atom) in the conduction band that excites the luminescent centers contained in the material with consequent emission of a low energy photon. The intensity of this signal is proportional to the absorption coefficient and will contain the XAS signal related to the environment of the absorbing atoms. The site selectivity between luminescent and nonluminescent zones under investigation can be achieved provided that the two families are sufficiently spatially separated (10 nm (Rogalev and Goulon 2002) for soft X-rays, an order of magnitude more for hard X-rays). In practical terms, the data collection is realized by collecting the optical emission from the sample with a lens focusing on an optic fiber leading to the entrance slit of a wavelength dispersive optical monochromator. Different portions of the emission spectrum can be thus used for the data collection permitting the selection of the luminescent sites.

Data Analysis Procedures It can be seen in Eq. 2 that each particular coordination shell contributes to the total signal with a sine oscillation in k space with frequency 2R (plus a term ϕ(k, Rj) weakly depending on k). Here we provide a description of the standard data analysis procedure (Lee et al. 1981): the raw absorption spectrum (Fig. 4a) the oscillating part χ is isolated by subtraction of a structureless atomic background (Fig. 4b). When the χ signal is Fourier transformed (FT), peaks appear in the spectrum each

F. d’Acapito

a

1.8

Absorption Coeff. (Arb. Units)

1240

1.6

b 3

1.4

2

1.2

1 k2 * x(k)

1 0.8 0.6 0.4 0.2

-3 -4

d

4

150

3

100

2.5

50

2 1.5

-100 -150

0 3

4 R [Å]

5

20

15

20

0

1

2

15

-50

0.5 1

10

200

3.5

0

5

k [Å 1]

q2 * x(q)

-3

0

-

Energy (eV)

Magnitude of the F T [Å ]

-1 -2

0 10500 11000 11500 12000 12500 13000 13500

c

0

6

7

8

-200

0

5

10 q [Å -1]

Fig. 4 Example of EXAFS data treatment relative to crystalline Ge. (a) raw absorption spectrum, (b) extracted EXAFS spectrum χ(k), (c) Fourier Transform, (d) Back Fourier Transform

corresponding to a different coordination shell (Fig. 4c). By windowing the FT in a way to leave only one peak and by applying to this function a back Fourier transform, the contribution from only one coordination shell is obtained (Fig. 4d). The basic principle of EXAFS data analysis is to reproduce the filtered experimental data (Fig. 4d) with a model based on the expression shown in Eq. 2. In that expression, two kinds of variables are present: • Parameters linked to the photoelectron interaction with the medium, like the backscattered amplitude A(k, Rj), phase φ(k, Rj), photoelectron mean free path λ, and S02. • Parameters linked to the local atomic structure, like the number of neighbors Nj, the bond length Rj, and the Debye-Waller factor σj. The former parameters can be calculated from ab initio methods. Several codes, like FEFF (Zabinsky et al. 1995; Ankudinov et al. 1998), GNXAS (Filipponi et al. 1995), and EXCURVE (Binsted) are available for this purpose. These functions can be used in a multiparameter fitting procedure with variable atomic structure parameters, to reproduce the experimental data. Different strategies of data fitting are used, the fit being done on Fourier-filtered EXAFS data (Lee et al. 1981) or directly on the FT (Zabinsky et al. 1995; Newille 2005). In both cases, the analysis can be done on a

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part of the spectrum (namely, the first shell) or on the whole. Other codes fit directly the absorption coefficient, making no use of Fourier Transform (Filipponi and DiCicco 1995), but in this case, a model accurately describing the entire structure needs to be considered. In general, bond lengths are determined with relatively high accuracy (0.02 Å, or better) whereas amplitude parameters (N and σ) are determined within 10%. A different approach is made on the XANES region. Here, since the calculation of the absorption coefficient is quite time consuming, the spectrum is reproduced at a qualitative level starting from a supposed three-dimensional model of the structure. Different codes based on a real space approach, like CONTINUUM (Tyson et al. 1992), FDMNES (Joly 2001), FEFF9 (Rehr et al. 2009) permit a modeling of the XANES part starting form a given structure and using the FMS approach. Other codes based on band structure methods (WIEN2K (Blaha et al. 1990), SPRKKR (Ebert 1998), XSPECTRA (Gougoussis et al. 2009)) are also available. Quite recently a new code called MXAN has been presented (Benfatto et al. 2001), permitting the structural refinement using the XANES part of the spectrum, with interesting results in biological applications (Della Longa et al. 2001).

Applications of XAS to Sol-Gel Materials A variety of examples on the use of XAS on materials obtained by sol-gel route will be presented here, with particular attention to materials for optical applications, catalysis, nanoparticles, and batteries.

Materials for Optics A considerable work has been carried out on Rare-Earth doped phosphors for applications in white Light Emitting Diodes. Indeed, these devices emit light in the blue or near ultraviolet (UV) portion of the spectrum and need a conversion agent to obtain white emission. Sol-Gel revealed to be a powerful technique for the realization of phosphors via treatments using low temperatures. Potdevin et al. (2010a, b) have studied the effect of Acetylacetone (acacH) as a chemical modifier in the synthesis of Tb-doped Y3Al5O12 (Yttrium Aluminum Garnet, YAG) phosphor in powder form. Indeed, sol-gel procedures revealed to be interesting with respect to the preparation via solid-state reaction involving high temperatures (T 1500 C) as this process can lead to byproducts spoiling the luminescent properties. In the procedure of production of YAG, it was found that using acacH instead of hydrolyzing the solution with water (Potdevin et al. 2010a), it was possible to obtain better materials using lower temperatures for treating the xerogels. A XAS investigation at the Y-K edge carried out on xerogels obtained with or without acacH revealed that in the former case the realization of a crystalline order around Y started already at 600 C whereas 800 C were necessary in the case of the materials obtained with the standard hydrolysis method (Fig. 5).

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F(R) /arb.units

1100 °C 800 °C 700 °C 600 °C 400 °C 200 °C xerogel 0

2

4

6

8

R/Å −1

k/A

F(R) /arb.units

1100 °C 800 °C 700 °C 600 °C 400 °C 200 °C xerogel 0

2

4

6

8

R/Å

Fig. 5 Comparison between the Fourier Transforms of XAS data at the Y-K edge in xerogels obtained with standard hydrolysis method (left) or with acetylacetone (right) (Figure Reproduced from Ref. (Potdevin et al. 2010a) with permission from The Royal Society of Chemistry). The peaks appearing around 3A are due to the second coordination (cationic) shell and indicate the formation of a crystalline nucleus. In the right panel (relative to samples prepared with acacH the second shell peak start to appear already at 600 C whereas the same peaks appear at 800 C in the case (left panel) of samples prepared with the standard hydrolysis method

The same behavior was evidenced in the environment of Tb dopants, confirming that the crystalline order was realized at lower temperatures in the matrix obtained with acacH and that dopants are incorporated in crystalline nuclei since the early stages of the process. Nakajima et al. (2013) have studied the incorporation of Ca-α-SiAlON phosphor doped with Eu (again for applications in white LEDs) in TiO2-SiO2 matrices

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obtained by sol-gel. Analyzing the XANES spectra, the authors have found that Ti was 5-coordinated for low Ti content (10 mol%) and turned predominantly 6-coordinated for high content (30 mol%). Eu contained in the phosphor was affected by the incorporation in the sol-gel matrix as the ratio Eu2+/Eu3+ was altered with a reducing effect (higher Eu2+ content) following the incorporation. Ye et al. (2013) have demonstrated the possibility of doping Mg2TiO4 with Mn via a sol-gel method with the aim of obtaining a red phosphor to be used in the wavelength conversion in white LEDs. In this case, X-ray Diffraction (XRD) revealed the formation of the Mg2TiO4 phase after annealing the xerogel at 1300 C. The chemical state and location of Mn dopants was elucidated by XAS at the Mn-K edge. By comparison with model compounds, the valence state of Mn was defined at Mn4+ and the detailed analysis of the XAS data (3 coordination shells) permitted to recognize MnTi as the incorporation site of this metal as required for the formation of the phosphor. Light conversion is needed not only in LED technology but also in solar cells. Indeed, the mismatch between the gap of Si and the solar spectrum leads to a limited efficiency of this kind of cells. To overcome this problem, particular coatings for solar cells have been proposed, called Quantum Cutting materials, capable to convert energy from the UV-visible (VIS) region into Near Infra-Red (NIR) where Si cells have a high efficiency. Sol-gel is an interesting method to produce these materials for the capability of covering large areas and the limited cost of the process. Terra et al. (2013) have investigated the formation of Tb3+-Yb3+ doped ZrO2 nanocrystals obtained by sol-gel. The two rare-earth ions have different roles: Tb3+ is the sensitizer that absorbs light in the VIS range and then transfers the excitation to Yb3+ (activator) which emits in the NIR. During the production of the material, it is of paramount importance that the dopants enter the matrix in the Zr-substitutional site and keeping the 3+ valence state. Materials were obtained with the one-step sol-gel method and contained different amounts of Yb (0–20 mol%) and fixed (1 mol%) amount of Tb. XRD showed that Yb favors the formation of cubic zirconia for content values above about 5 mol%. XAS at the Tb-L3 edge elucidated the valence state of Tb in the xerogel and calcined materials. In the former case, Tb is 3+ for any value of Yb content, presumably due to the fact that Tb nitrate was used as precursor. Upon calcination, formation of Tb4+ was observed for low-Yb containing materials, whereas Tb resulted to keep the 3+ state for Yb at 20 mol%. The inhibition of the unwanted oxidation of Tb was explained by the effect of oxygen vacancies generated by Yb (aliovalent with respect to Zr) and permitted to define the optimum conditions for obtaining an effective QC material. Similar effects linked to the valence state of rare earths have been also reported in Carvalho et al. (2016) where the valence state of Pr in Gd-doped ZrO2 obtained by a standard sol-gel procedure was controlled via the Gd content as revealed by XAS at the Pr-L3 edge. The authors have shown that with Gd at 10 mol% the structure of ZrO2 is cubic and that Pr is exclusively in the Pr3+ state. Codoping can be used also to improve the behavior of photocatalysts as shown in the case of TiO2 by Majeed et al. (2015). TiO2 is known to be an excellent photocatalyst but it suffers from the fact that its gap is 3.15 eV so it can use only

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about 3–5% of the solar spectrum. Doping TiO2 with metals can contribute to an improvement of its properties as in the case of Mo doping. TiO2 was prepared via sol-gel with the addition of molybdic acid as Mo precursor. Mo was added at different concentration values from 1% to 10% and the obtained samples were analyzed by XAS at the Mo-K edge. The XAS data revealed that Mo is present in two phases: substitutional phase in TiO2 and a separated MoO3 phase. Mo is totally substitutional for the samples at 1% and the amount of this phase decays with Mo content, being only 7% in the highly doped sample. Tests of efficiency were carried out measuring the photodegradation of Methylene-Blue pigment in presence of the catalyst under visible irradiation. The result was that the better performing sample was that containing 5% Mo out of which 30% is in substitutional phase, demonstrating that the interplay of all three components (TiO2, Mo: TiO2, and MoO3) contribute to the improvement of the efficiency of the catalyst. Studying luminescent materials permits the exploitation of nonconventional XAS techniques like X-ray Excited Optical Luminescence (XEOL). This technique has been used to characterize Ge nanoparticles (NP) (Little et al. 2014) in free-standing form (comparing in this case oxygen-terminated particles with hydrogen-terminated particles) and Ge nanoparticles embedded in silica produced by the sol-gel route. The spectra at the Ge-K edge were collected in transmission mode (sensitive to all the Ge species) and XEOL mode (sensitive to Ge species in luminescent regions) for comparison. In the case of Ge nanoparticles surrounded by O species (O-terminated free standing NPs or NPs embedded in silica), both spectra revealed a predominance of Ge-O bonds meaning that the luminescence comes from the oxidized zones. In the case of free NPs terminated with hydrogen, the XEOL spectrum exhibited a higher disorder and a comparison with Molecular Dynamics simulations permitted to establish that the luminescent regions were the most disordered one at the boundaries of the particle. XEOL detection was also used to characterize Rare-Earth doped nanospheres for medical applications (Fortes et al. 2014). Hollow and bulk silica particles coated with Er2O3 and Yb2O3 were characterized by XAS in fluorescence and optically excited mode. The comparison between the spectra obtained in the two ways showed no differences meaning that all the rare-earth species were present in luminescent portions of the samples. Further uses of XAS for the characterization of materials for optical applications produced by sol-gel can be found in the case of Er-doped Silica-Hafnia waveguides for optical amplifiers (d’Acapito et al. 2008; Afify et al. 2007), glass ceramics (Van et al. 2015), and scintillator for detectors (Liu et al. 2015b).

Nanoparticles The sol-gel process permits to produce nanocrystalline materials with a good control of composition and size that reveals to be invaluable for a large variety of systems. Metal oxide nanoparticles, namely, find applications in several fields like gas sensors, catalysts, electrodes for batteries and transparent electrodes.

40

X-ray Absorption Spectroscopy Studies on Materials Obtained by the. . .

1245

Caetano et al. (2014) have studied in detail the structural evolution from the gel to the densified material of SnO2 nanoparticles using time-resolved XAS at the Sn-K edge. The synthesis was based on Tin Chloride (SnCl4 • 5H2O) dissolved in ethanol that was subsequently hydrolyzed with water. Quick-XAS data were collected at a frequency of 2 Hz during hydrolyzation and XAS, and Raman spectra were collected at the same time in a specially conceived cell. The first 30 min of the process have been described as a progressive loss of chlorine ligands for Sn substituted by oxygen ligands. The XAS data as a function of time could be reproduced as a linear combination of different complexes in solution [SnClx(H2O)6x]4x (with x = 3, 4 and 5) and the spectrum of SnO2 nanoparticles. With the support of Raman data, it was possible to describe the evolution of the various complexes involving Sn as shown in Fig. 6. Time-resolved XAS was also used in the description of the growth of TiO2 nanoparticles in sol-gel as reported in Stotzel et al. (2010). In an ad hoc conceived reaction, cell Titanium tetraisopropoxide was solved in isopropanol and successively a water-isopropanol solution was added to carry out the hydrolysis of the solution. At the same time, XAS spectra were collected with a duration time of 5 s (10 successive spectra were averaged to improve the signal to noise ratio) up a total process time of 1500–2000s. The analysis of the XAS data was carried out by linear combination of the spectra at different times with model compounds. The formation of TiO2 particles involves different steps: polynuclear titanium species are created and,

Fig. 6 Evolution of different species [SnClx(H2O)6x]4x and SnO2 in solution during hydrolyzation and annealing. Results from XAS and Raman data collected in situ at the same time (Reprinted with permission from Ref. (Caetano et al. 2014). Copyright 2014 American Chemical Society)

1246

F. d’Acapito

after a few minutes, oligomers of the type Ti11O13 and Ti12O16. The consumption of the precursor is particularly evident as it exhibits a particularly strong pre-edge resonance at about 4970 eV that is considerably damped in the oligomers and TiO2 species due to the different site symmetry in the two cases. In an intermediate phase, the oligomers start to form titania nanoparticles and, at the end, they are totally converted in TiO2. The doping of SnO2 nanoparticles with Sb, obtained by sol-gel processing, has been studied by V. Geraldo et al. (2010). In this case, gels at different Sb content (3–16%) were prepared from SbF3 and (SnCl4 • 5H2O) precursors. The obtained xerogels were treated at 200 C or 500 C to study the evolution of the valence state and incorporation site of Sb. It was evidenced that after treating at 200 C, Sb is pentavalent if present at low concentration ( 4%) and a trivalent fraction is present for higher Sb content. After firing at 500 C, Sb is predominantly pentavalent. Sb5+ species enter the SnO2 matrix in a substitutional site. Conversely Sb3+ species, more abundant at high concentration values, remain grafted at the surface of SnO2 particles. The formation of Cooxide nanoparticles in porous silica has been studied by Liu et al. (2015a). In this case, Cobalt doped silica at different Co/Si molar ratio (5%, 10%, 25%) was obtained by sol-gel using Tetraethoxysilane (TEOS) for silicon and [Co(NO3) • 6 H2O] as precursor for Cobalt. The sols were first dried at 60 C and then ground in fine powder and calcined at 630 C. The materials were successively tested with a hydrothermal treatment (HT) consisting in annealing the samples at 550 C for 40 h in an atmosphere containing water vapor at 75 mol%. XAS measurements at the Co-K edge were carried out to reveal the chemical status of the metal prior or after the HT. All samples revealed to be stable under HT; no remarkable differences were present in the XAS spectra with respect to the non-HT ones. In samples containing 5% Co, Cobalt was found tetracoordinated to oxygen with a faint signal from a Co second shell. On the other samples with higher Co content, a well-defined second shell signal was visible that was attributed to the formation of Co3O4 nanoparticles. Again on the topic of metallic nanoparticles, Takao et al. (2012) have demonstrated the production of metallic Pd and PdO particles in hollow silica nanoparticles. This method possesses the advantage of producing metallic clusters with an extremely reduced number of atoms (4–60) that is difficult to be obtained in other ways. The hollow nanoparticles were produced starting from a spherical template of Pd12L24 where L is a suitable ligand (triethoxysil group). The silica is deposited on the surface of this template by using tetramethoxysilane with the standard procedure. Successive calcination at 400 C in air eliminated the ligands leaving small clusters with 12 Pd atoms inside the hollows of the spheres. XAS at the Pd-K edge showed that the clusters are made of PdO and that further treatment with H2 leads to the formation of metallic Pd. Further studies on nanoparticles obtained by sol-gel and characterized by XAS have been carried out for the production of magnetic semiconductors (Bilecka et al. 2011; Kumar et al. 2014; Hu et al. 2011; Liu et al. 2010) and catalysts (Santos et al. 2012; Prieto et al. 2010; Mathew et al. 2011).

40

X-ray Absorption Spectroscopy Studies on Materials Obtained by the. . .

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Materials for Catalysts and Sensors Metallic nanoparticles (MNP) embedded in a porous material are largely used in heterogeneous catalysis but these systems easily undergo deactivation by sintering. A way to limit this problem is to host them in a suitable host to maintain their dispersion. Among others, silica Aerogels have been proposed but their use remains limited due to their high production costs. Recently, Kristiansen et al. (2011, 2012) have succeeded in incorporating Cu in silica aerogels produced by Ambient Pressure Drying, a method that greatly simplifies the production of this kind of material. In the first paper (Kristiansen et al. 2011), the group has demonstrated the dispersion of Cu species in the aerogel by looking to the XAS spectra collected at the Cu-K edge and verified that no Cu-Cu bonds were present. In a subsequent study (Kristiansen et al. 2012), the group has used XAS to study the formation of Cu nanoparticles in presence of a reductive atmosphere (5% H2 in He) and at high temperature. The XAS spectra presented metal-metal bonds with a reduction of the first shell coordination number due to the reduced size of the particles (estimated to contain about 80 atoms). Data were collected at increasing temperature values and differences in the spectra were evidenced suggesting a change in shape and local symmetry of the clusters. In order to interpret the XAS experimental data, XAS spectra were simulated based on the results of atomistic simulations (in the framework of the DFT theory) of clusters of different shape. The results of the study are collected in Fig. 7. At about 300 C, Cu forms nanoclusters that, from the number of Cu first neighbors observed (8) and comparison with DFT theoretical data, are identified as clusters with Td symmetry (fcc structure) containing about 80 atoms. The size of these particles (about 1.4 nm) is well matched with that of the typical pores ( 0.345)

~180–300 (depends on composition) 400–1500 (depends on composition)

23,000 123 (e33/e0)

20–35(Pr, depends on composition)

4 mm, or 3 m depends on composition mm2 m

mm2

200–1500 (e33/ ε0 depends on composition) ~3800 (la/Zr/Ti = 8/ 65/35)

20–55 (Pr, depends on composition)

4 mm (x < 0.53) 3 m (x > 0.53)

All data in the table refer to bulk material, except where indicated.

PLZT La-doped Pb(ZrxTi1–x)O3 SbSI Bi4Ti3O12

450–250 (depends on Zr/Ti ratio) 90–300 depends on composition 20 675

PZT Pb(ZrxTi1–x) O3

2.6  101

1.7  101 (la/Zr/ Ti = 8/65/35)

Depends on composition

30 (d31) 65(d33) (calculated, x = 0.75, film)

23 (d33)

1300 (d33) 20 (d33)

~20–40

71 (d31)

57 Ferroelectric and Piezoelectric Properties 1683

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Y. Xu and J. D. Mackenzie

ð Usc  ð1=CÞ I s dt ¼ ðA=CÞ½Δ PÞTT 21

(16)

The pyroelectric current is obtained from the slope of the curve Us(T ). (d) Digital integration measurement (Li et al. 1984): A schematic diagram of the experimental setup for the digital integration method for accurately measuring the pyroelectric coefficient at room temperature is shown in Fig. 11. The sample S, in the form of a thin slab, is enclosed in a protective tube inserted in a water bath along with a (Beckmann) thermometer. The initial temperature T1 is measured when the system is in thermal equilibrium. The initial temperature T1 must be maintained so long as there is no pyroelectric charge in the sample. The water bath is then heated to a temperature T2; this temperature must be maintained for a long period of time. Meanwhile, the pyroelectric charges are integrated by an integrating circuit, to produce a positive (or negative) integrated voltage. When the integrated voltage reaches a certain absolute value, the integrating circuit gives a discharge pulse. A control circuit controls the time interval of the discharge time with a quartz oscillator, while a positive (or negative) pulse counter records the discharge time. During the period of measurement, the voltage between the electrodes of the sample is less than several microvoltages, thus, the condition of zero electric field is essentially maintained. Assumed stable data recorded at the positive and negative counters are N+ and N, respectively; the total pyroelectric charge is then ΔQ ¼ δQðN þ  N  Þ ¼ δQ ΔN

(17)

Fig. 11 Block diagram of the experimental setup of the digital integration method, for measurement of the pyroelectric coefficient (Li et al. 1984, with permission)

57

Ferroelectric and Piezoelectric Properties

1685

where δQ is the discharge per pulse. The pyroelectric coefficient is thus given by: p ¼ ΔP=ΔT ¼ ðΔQ=ΔT ÞA ¼ ðδQΔN Þ=ðT 2  T 1 ÞA

(18)

For common pyroelectric materials, when the electrode area A is 1 cm2 and Δ T is about 1
C, the value of ΔN ensures results with an accuracy of 4 to 5 digits and fairly high reproducibility.

Piezoelectric Properties Piezoelectric Eqs. P. Curie and J. Curie discovered the piezoelectric effect in 1880. It was found that, when a compressive or a tensile force was applied on some crystals along some special directions (e.g., a quartz) electrical charges could be created on the corresponding surfaces of the crystal and the size of the created charge was proportional to the strength of the applied force. This phenomenon is called the “piezoelectric effect.” All ferroelectric crystals show a piezoelectric effect. The piezoelectric effect can be described by piezoelectric equations. On the basis of thermodynamic principles, piezoelectric equations can be derived (e.g., see Xu 1991). These equations describe linear relationships between the four variables: stress tensor [T], strain tensor [S], electrical field vector E, and electric displacement vector D. The piezoelectric equations can be expressed as four kinds of equations, depended on the variables. Selecting E and T as variables, we have: Si ¼ ΣsEij T j þ Σdin En and Dm ¼ Σdmi T i þ ΣeTmn En

ði, j ¼ 1, 2, . . . , 6; and m, n ¼ 1, 2, 3Þ

(19)

These are called the first piezoelectric equations, sij are elastic compliance coefficients, emn are dielectric constants, and dmi are piezoelectric coefficients (or piezoelectric moduli). If we taking E and [S] as variables, the second piezoelectric equations will be as follows: T j ¼ ΣcEij Si  Σejn En ði, j ¼ 1, 2, . . . , 6; and m, n ¼ 1, 2, 3Þ Dm ¼ Σemi Si þ ΣeTmn En

(20)

where cij are elastic stiffness constants and enj are piezoelectric constants. The third and the fourth piezoelectric equations are T j ¼ ΣcD jj Sj  Σhjn Dn En ¼ hnj Sj þ βSn3 D3

(21)

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Y. Xu and J. D. Mackenzie

and Si ¼ ΣsD ij T j þ Σgim Dm , En ¼ Σgnj T j þ ΣβTnm Dm ,

(22)

where βnm are the dielectric impermeabilities, and both hnj and gnj are piezoelectric constants. The boundary conditions for these piezoelectric equations are important: (a) The condition “mechanically free” stipulates specifically that boundaries of a piezoelectric sample (e.g., a piezoelectric vibrator) can move, i.e., the vibrator vibrates with a variable strain and zero (or constant) stress. Under this condition, the coefficients in these equations carry a superscript “T”; e.g., eTmn is the dielectric constant at constant stress. (b) The condition “mechanically clamped” stipulates specifically that the boundaries of a vibrator cannot move. This condition means that, when the frequency of the applied voltage is much higher than the resonance frequency of the vibrator, the strain is constant (or zero), while the stress varies. In this case, the coefficients in these equations carry a superscript “S”; e.g., eSmn is the dielectric constant at constant strain. (c) The condition of “electrical short circuit” implies specifically that the electric field E = 0 (or a constant), while the electric displacement D 6¼ 0 inside the vibrator. This is the case when the two electrodes on the surface of the crystal sample are electrically connected (or the electric potential on the entire surface of the sample is constant). Under this condition, the coefficients in these equations carry a superscript “E”; e.g., sEij (or cEij ) is the elastic compliance (or stiffness) coefficient at constant electric field. (d) The condition of “electrical open circuit” corresponds to the case when all the free charges are kept on the electrodes of the sample (electrically insulated) and the internal electric field E 6¼ 0, while D = 0 in the sample. In this case, the coefficients in these equations carry a D superscript “D”; e.g., sD ij (or cij ) is the elastic compliance (or stiffness) coefficient at constant polarization. The relationships between the four piezoelectric constants, dnj, hnj, gnj, and enj, are dnj ¼ ΣeTnm gmj ¼ Σeni sEij ðunit is C=N or m=VÞ,   enj ¼ ΣeSnm hmj ¼ Σdni cEij unit is C=m2 or Vm=N ,   unit is m2 =C or N=Vm , gnj ¼ ΣβSnm d mj ¼ Σhni sD ij

(23)

hnj ¼ ΣβSnm emj ¼ Σgni cD ij ðunit is N=C or V=mÞ, According to Neumann’s principle, physical properties in an anisotropic crystal are also anisotropic and they can be described by tensors of different ranks, such as the second-rank tensor emn, the third-rank tensor dnj, and the fourth-rank tensor sij, etc. The piezoelectric constant tensor is a symmetric third-rank tensor, therefore, its 27 components can be reduced to 18 independent components in the abridged two-subscript notation. For various point groups, the number of independent components of tensor physical properties can be further reduced.

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Ferroelectric and Piezoelectric Properties

1687

Electromechanical Coupling Factor Besides the four piezoelectric constant, another parameter, the electromechanical coupling factor k, is a very useful parameter for a piezoelectric material (IRE Standards 1961). The coupling factor k represents the efficiency of energy transfer between electrical energy and mechanical energy. If an electrical signal (or a mechanical force signal) is applied to a piezoelectric vibrator, an energy transfer occurs due to the piezoelectric effect, or counter-piezoelectric effect. The definition of the coupling factor, k, is. k2 = [electrical energy output]/[mechanical energy input] (by piezoelectric effect). or k2 = [mechanical energy output]/[electrical energy input]. (by counter-piezoelectric effect). In general, it can be derived that: ðkni Þ2 ¼ ðd in Þ2 =eTnn sEii ði ¼ 1, 2, . . . , 6; and n ¼ 1, 2, 3Þ

(24)

Piezoelectric Vibrators and Determination of Piezoelectric and Elastic Parameters by the Transportation Line Method In general, elastic coefficients, piezoelectric constants, and electromechanical coupling factors are determined by the transportation line method, which was recommended as IRE standard (1957, 1958, 1961). This method is a dynamic method, because the measurement sample is made as a vibrator and measured in a vibration state at a frequency near its intrinsic resonance frequency. When a small AC voltage is applied to a piezoelectric sample (poled single crystal, or poled bulk ceramic) having a pair of electrodes on opposite surfaces, the sample vibrates at a frequency near its resonance frequency, fr . Figure 12 shows a sample inserted in a transportation network. By changing the frequency of the signal generator, a resonance frequency, fr, and an anti-frequency, fa, are measured. The two frequencies correspond approximately to the maximum and minimum of the output voltage through the network, respectively. Applicable circuits of transportation line include Π-type networks, Γtype networks, frequency scanning analyzers, and other special circuits. As shown in Fig. 12, a circuit of a conventional Π-type network for the constant voltage method includes a transmission network, a signal generator, a high frequency voltmeter, and a frequency meter. In this network, the components arranged at symmetric positions have the same magnitude. Ri and RT are noninductive resistors, the latter being variable. It is required that Ri  RT and RT  |Zm|; |Zm| is the impedance of the sample. As the frequency of the signal generator changes, the frequencies at maximum transmission and at minimum transmission are measured with the aid of the high-frequency voltmeter and the frequency meter. The two frequencies are approximately equal to the frequencies fr and fa of the piezoelectric samples.

1688

Y. Xu and J. D. Mackenzie

Fig. 12 A transport at mission network for measurement of resonance frequency and antiresonance frequency of a vibrator sample (from IRE Standard 1961, with permission)

In general, applying a variable frequency AC voltage to a piezoelectric sample, which has a specific shape and a specific polarization orientation, may excite a specific resonance mode. Thus, fr and fa of the sample can be recorded. By using fr, we obtain a related “frequency constant” N = frl, where l is the sample’s dimension along the direction of major vibration. The value of fr and N depend on the vibration mode of the sample. Based on the vibration equation of the piezoelectric vibrator, the electromechanical coupling factor, k can be calculated: k2  π ðf a  f r Þ=2fr:

(25a)

By using the data of fr and fa, the components of the elastic coefficient and the piezoelectric constant can be calculated from the dynamic equation of the piezoelectric vibrator and the piezoelectric equations. For example, a piezoelectric sample is a bar along the x-axis, and an AC voltage is applied on the sample along the z-axis, which is the poling direction of the sample. A “longitudinal extension mode” vibration (recording the mode as “31”) of the sample is excited by the applied signal. From fr31 and fa31, we can calculate N31 and k31: ðk31 Þ2  π ðfa31  fr31 Þ=2fr31

(25b)

From the vibration equation:   1=2  2  =2l or ¼ cE11 =ρ =2l : fr ¼ ρsE11 Recalling Eq. (17), we have

(26)

57

Ferroelectric and Piezoelectric Properties

ðk31 Þ2 ¼ ðd31 Þ2 =eT33 sE11

1689

(27)

Therefore, we obtain the data k31, d31, and sE11 from the measurement of fr and fa in vibration mode of “31” and eT33 is measured independently. After a piezoelectric ceramic is poled by a DC field higher than the coercive field Ec, the sample has symmetry of 1mm. For this piezoelectric ceramic, a typical process chart for the determination of dielectric constant, piezoelectric constant, elastic coefficient, and electromechanical coupling factor is shown in Fig. 13.

Measurement of Piezoelectric Constants by a Static (or Quasi-Static) Method In the static method for measure of piezoelectric constants, the sample is not in a vibration state during the measurement. If only a mechanical stress is applied on the measured sample without an electric field, from Eq. (12) we have: Dm ¼ d mj T j ðm ¼ 1, 2, 3; and j ¼ 1, 2, . . . , 6Þ When a static measurement is being done, an external force with a known magnitude, F3, is applied to the sample in a fixed direction; consequently, charges accumulate at the sample’s surfaces due to the piezoelectric effect. These accumulated surface charges have a magnitude (per unit area) exactly equal to the electric displacement per unit volume. If the direction of the applied stress is in the direction of the sample’s polarization axis, only the T3 component of the stress exists and D3 ¼ d33 T 33: If the total force corresponding to the component T3 is F3, the total charge on the surface electrode is Q = d33F3. Suppose C is the static capacity of a sample having a charge Q; a voltage V = Q/C will be built up across the two electrodes of the sample. Then we have: D33 ¼ Q=F3 ¼ CV=F3 :

(28)

A spring, a lever, or some weight can apply the force F3 on the sample. An electrostatic voltmeter and a charge-integrating meter can measure V and Q, respectively; thus, d33 can be calculated. In practice, the technique of stress relief is usually used instead of a pressure application. There are other techniques also available for measuring piezoelectric constants. Using some special commercial equipments with digital display, the process of measurement is more convenient and prompt. For example, the well-known “Berlincourt d33 meter” or “ZJ-2 quasi-static d33 meter” is based on a quasi-static condition. The principle of quasi-static method is as follows. An alternating pressure

1690

Y. Xu and J. D. Mackenzie (1) Disk-extension vibration Co

E33T

s Fr

s11E

s13E

s12E d31

kp

(2) Plate-thickness longitudinal vibration Fr

c33D c33E

kt

(3) Bar-longitudinal vibration (terminal faces electrodes) Fr

s33D

s33E

k33

d33

(4) Thickness shear vibration Fr

c44D c44E

k15 Co

E11T

s44E d15

Fig. 13 A typical process chart for the determination of dielectric, elastic, and piezoelectric constants (from IRE Standard, 1961, with permission)

of low frequency (several Hz, or a few hundred Hz) is applied both on a piezoelectric sample and on a reference, which has a known piezoelectric constant d33, while the created charges in both sample surfaces are collected separately. By comparing the created charge in the sample with that in the reference, the value of d33 is computed and displayed. We must point out here that piezoelectric constants obtained by a static method are usually different from those obtained by a dynamic method, owing to the fact that the boundary conditions in the two methods are different.

57

Ferroelectric and Piezoelectric Properties

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Effective Piezoelectric Coefficient of a Piezoelectric Film/Substrate System Usually, a ferroelectric thin film sample fabricated by the sol-gel method is coated on a substrate, which is much thicker than the film. The piezoelectric constant, d33, of the thin film sample can be measured by the static (or quasi-static) method and the electromechanical coupling factor k33 can be calculated from the known dielectric constant eT33 and the elastic coefficient sE33. It is hard to measure a thin film/substrate system by a dynamic method, e.g., the transportation line method, because the vibrator is a thin film/substrate system and the major vibration is caused by the substrate, not by the film. A quasi-static method has been applied to measure a sol-gel fabricated PLZT (lanthanum-doped PZT, 400 nm thick) thin film on Pt/Ti/LTO (silicon dioxide) substrate. This technique measures the tip displacement of a cantilever sample, by using a laser interferometer, to obtain an “effective piezoelectric coefficient,” d31(eff) (Cheng et al. 1999). A piezoelectric cantilever was made and driven by a voltage V. In general, if the substrate has the same elastic modulus as the piezoelectric film and both are of the same thickness, t, and the same length, l, the tip displacement, δ, of the cantilever, for a driving voltage, V, is (Hoffmann et al. 2002):  δ ¼ V 3d31 l2 =2t2

(29)

where d31 is the piezoelectric constant of the film. This is a simplified calculation. However, calculation of a practical micro-cantilever is more complicated than that of an ideal cantilever. A tested sample was fabricated as a cantilever, 15 μm in width and 90 μm in length. The bottom electrode was a Pt/Ti layer on the substrate (LTO) and an upper electrode was Pt, coated on a PLZT surface. When a DC voltage was applied across the PLZT layer of the cantilever sample, the PLZT layer was shrunk due to the counter-piezoelectric effect and the cantilever was bent up, as shown in Fig. 14. The displacement of the cantilever tip can be measured by using a laser interferometer (e.g., Pan and Cross 1989). The displacement, δ, of a cantilever is approximately proportional to the applied voltage V, when V is not very high (V < 10 V). The relationship between δ and V can be expressed by the following equation:  δ ¼ V 3d31ðeff Þ s1 s2 t1 ðt1 þ t2 Þl2 =s21 t42 þ 4s1 s2 t1 t32 þ 6s1 s2 t21 t32 þ 4s1 s2 t2 t32 þ s2 t41 (30) where t1 and t2 are the thickness of the LTO layer and the PLZT films (t1 = 300 nm and t2 = 400 nm); s1 and s2 are the elastic compliance sE11 of the LTO and PLZT films, respectively; l is the length of the cantilever; and d31(eff) is the effective piezoelectric coefficient of the PLZT film (in fact, a Pt/PLZT/Pt/Ti/LTO system). A measured result is plotted in Fig. 15. From the linear part (10,000

10 < T608 < 65 (1 V/+1 V)

6

G/FTO/TiO2/Chrom 1/electrolyte: LIClO4, ferrocene in γ-butyrolactone/FTO/G

– 10 nm) form thicker films with a fraction of large pores or a tendency to macroscopic cracking after drying (Castricum et al. 2008d). Acid-catalyzed organosilica sols suitable for membrane formation usually have a radius of gyration Rg < 10 nm and Di = 1.2–1.6. The dominant growth mechanism under most typical solgel preparation conditions is diffusion-limited cluster aggregation (DLCA), which ultimately results in nanostructures with Di = 1.8–1.9 (Boffa et al. 2009; Lin et al. 1989; Maene et al. 1998). In DLCA the approach of particles is controlled by Brownian transport, and the rate of reaction is much higher than the rate of transport. A study on sols with -CH2-, -C2H4-, -C8H16-, p-phenylene (-p-C6H4-), and di-p-phenylene (-p-C6H4-pC6H4-) bridging groups showed that the values of Di are similar (Di = 1.5–1.6) when the radii of gyration are similar (Rg = 1.5–2.3 nm), even though the reactivity of the precursors varies (Castricum et al. 2011). Defect-free membranes with molecular sieving properties could be formed from all five sols. The acid concentration also has a profound influence on the formation of the pore structure of hybrid organosilicas. The evolution of nanostructure is very different depending on the [H+]/[Si] ratio (Castricum et al. 2014). In situ SAXS experiments on drying BTESE films with [Si] = 0.9 M, [H2O]/[Si] = 1, and [H+]/[Si] = 0.01 at room temperature showed a decreasing scattering intensity, while Di dropped to zero after 4000 s. This implies the formation of a homogeneous film without pores visible on the length scale of the scattering experiment (0.5 nm and larger). When the [H+]/[Si] ratio was increased to 0.1, a smaller decrease in scattering intensity was observed upon drying. Di decreased from 1.2 to 0.4 in the final as-dried film after 4000 s. This value is illustrative for a densifying film with a final “fractal-like” pore structure. A third film with [H+]/[Si] = 1 showed complex drying behavior. It developed initially in a similar way as the films with lower acid ratio, but then Di increased to 2.15 after 4500 s, followed by a gradual decrease to a final value of 1.5. While the detailed interpretation of this process is complicated, the final

103

Hybrid Materials for Molecular Sieves

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decrease suggests a slow densification process, however less than in the other two cases. What this study illustrates is that the presence of H+ is necessary to construct and/or maintain a porous network structure in the drying film (gel) during solvent evaporation. The naturally occurring capillary forces and compressive stresses during drying and film shrinkage promote the densification of the hybrid silica network into a dense structure with small or no pores, but these forces seem to be counteracted at high acid concentrations. Possibly the network is also strengthened by ongoing condensation catalyzed by H+ and/or by the positive charge on the polymeric colloidal network that may somehow stabilize it. Thermal processing of these BTESE solgels in N2 atmosphere at 523 K yielded powders with N2 BET surface areas of 0, 632, and 970 m2/g, respectively, and corresponding pore sizes < 0.3 (N2 size), 1.5, and 2.7 nm (Castricum et al. 2014). Hence, the presence of H+ during reaction and film drying leads to a more open network with larger pores. This study clearly demonstrates the importance of the H+ concentration to tailor the pore size of a hybrid silica membrane. Interestingly, the pore structure can be adapted in this way, i.e., by adjustment of the acid-to-Si ratio before the onset of physical drying, while the original rheological properties needed for film coating are retained.

Pore Structure in Thermally Consolidated Materials The presence of organic groups in the hybrid matrices after thermal treatment can be determined by Fourier transform infrared (FTIR) spectroscopy. The symmetric and asymmetric stretching of the organic groups can be found in the high-frequency range, i.e., 2800–3000 cm1 for the alkylene systems and 2950–3100 cm1 for the aromatic systems. Ethylene- and octylene-bridged materials show a broad band at 2800–2950 cm1, while the -CH2- bridged system has an additional pronounced peak at 2975 cm1, at similar position as the terminal methyl group from MTES. The pore structure of hybrid organosilica after thermal treatment differs significantly from the pore structure of microporous SiO2. Molecular dynamic (MD) simulations of -C2H4- and -C2H2- bridged systems suggest that the pore size distribution of hybrids is slightly shifted toward larger pore sizes than in SiO2 (Kanezashi et al. 2010; Shimoyama et al. 2013). Simulated pore size distributions of BTESE hybrid silica and SiO2 silica are shown in Fig. 3 (Chang et al. 2010, 2011). It is noted that the smallest gases and vapors relevant for separation with membrane process are H2O and NH3, both with a kinetic diameter 0.4 nm. This larger pore size is related to the organic bridging groups that make the Si to Si distances, i.e., the branching points in the 3D network, larger. The presence of a small fraction of large pores also explains the comparatively low permselectivities obtained in the separation of elementary gases (H2, CH4, CO, CO2), as discussed in more detail below. Much higher permselectivities can be

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Fig. 3 Simulated cavity size distribution of the pure silica and hybrid BTESE silica membranes (Reproduced from Chang et al. (2010) with permission of The Royal Society of Chemistry)

accomplished for larger gases and vapors such as propene (0.46 nm)–propane (0.50 nm) separation (Kanezashi et al. 2012a) and the separation of water from alcohols and other organic solvents, i.e., organic molecules with typical molecular diameters > 0.4 nm. Experimentally, BTESE powder made via a similar solgel process as BTESE membranes has an N2 BET surface area of 130–310 m2/g at 77 K and surface areas of 510–550 m2/g for ethene (C2H4) and acetylene (C2H2) at 273 K (Castricum et al. 2008a, 2011). The N2 adsorption isotherms are of type I and without hysteresis for both materials indicating full microporosity (Castricum et al. 2011). The size of N2 in BET experiments is assumed to be 0.30 nm, while the size of C2H2 is only 0.24 nm. So the latter value reflects the large fraction of ultra-small pores present in BTESE. BTESM has slightly smaller pores than BTESE owing to its shorter carbon bridge, i.e., almost no N2 adsorption at 77 K, and somewhat lower values for C2H4 and C2H2 than BTESE. The pore structure and pore connectivity can be strongly altered by co-condensation of other silicon alkoxide precursors such as MTES. Mixed BTESE–MTES powders (1:1 molar ratio) were found to have much smaller pores than BTESE. No adsorption of N2 occurred, while the adsorption of C2H2 indicated a surface area of 1340 m2/g (Castricum et al. 2008d). Moreover, the skeletal density of BTESE–MTES was 10% smaller than that of BTESE (1.5 vs. 1.7 g/mL; cf. dense SiO2 with density of 2.2 g/mL). Hence, the introduction of terminal methyl groups in the hybrid structure seems to promote the overall porosity and the fraction of ultrasmall (0.3 nm are present but only accessible to CO2 (340 m2/g), not to N2. In comparison, MTES-derived methylated silica with the same overall CHx:Si ratio and microporous SiO2 silica both showed very little N2

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adsorption (12 and 33 m2/g, respectively), and their adsorption capacities for C2H2 were also smaller (211 and 261 m2/g, respectively). So the average pore size of BTESE–MTES hybrid silica is smaller than that of MTES- and TEOS-derived silica. TEOS–MTES hybrids (1:1 molar ratio) even seem to have larger pores than the other systems, with a N2 BET surface area of at least 400 m2/g. Materials with longer organic bridging groups, i.e., octylene, phenylene, and biphenylene, also showed little or no uptake of N2, similar to BTESE and BTESM materials (Castricum et al. 2011). Despite their lengths, their pore sizes are not much larger than those of BTESM or BTESE. The skeletal densities are typically lower (1.2–1.4 g/mL) than the densities of the shorter-chain analogues owing to their higher organic contents. It is noted that the helium (0.20 nm) pycnometry measurement used for skeletal density determination required long equilibration times for all hybrids, which is suggestive of very small pores and relatively poor interpore connectivity. The smallest micropores were found in the octylene-bridged sample. They were estimated to be even smaller than in BTESM. Also the lower adsorption potential for CO2 compared to all other investigated materials suggests that in the octylene-bridged material, gas adsorption is hindered, probably by the flexibility and/or conformation of the long alkylene bridge.

Supported Membrane Preparation and Determination of Pore Size Distribution Membrane Preparation The small pore size and relatively low porosity of hybrid organosilicas after thermal treatment finally results in a material that may have good intrinsic molecular sieving properties but that is also very resistive to transport of molecules. It is therefore imperative to keep the thickness of the separating layer minimal. For this reason, hybrid organosilica membranes are always used as thin films supported by a macroporous support structure that maintains mechanical integrity. Organosilica sols are made by acid-catalyzed solgel processing of bridged silsesquioxane precursor molecules in an alcoholic solvent, usually ethanol. The bridged precursor is present at a concentration of 0.9–1 M. A typical reaction in ethanol involves [Si] = 1.8–2.0 M (two Si centers per bridged precursor), [H+]/ [Si] = 0.1, and [H2O]/[Si] = 2–3. The pH of the resulting solution is < 1 (Castricum et al. 2008b; Kreiter et al. 2011). To carry out the reaction, distilled water and nitric acid are added to dry ethanol, and the precursor is then added under vigorous stirring. The fast occurring hydrolysis reaction is very exothermic and may lead to a sudden initial temperature increase. Although many variations of the solgel process including room temperature processing are possible, the reaction is often carried out in a water bath set at 60  C for 1–3 h under continuous stirring. The resulting stock sol can be stored for later use, although it should be kept in mind that the sol is not thermodynamically stable. These sols show noticeable growth at 18  C over

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a period of weeks, and we even noticed slow growth of hybrid silica sols at 78  C over a period of a few months. This sets a limit to the lifetime of the sol. In most studies, the hybrid separation layer is dipcoated onto an α-aluminasupported γ-alumina substrate with a single or double dipcoating step. Just before the hybrid silica sol is applied to form a membrane film, it is diluted with ethanol by a factor of 5–25 on the basis of volume to obtain a fast-drying film with good film formation properties and low (Newtonian) viscosity for the dipcoating process. After drying, the membranes are heat-treated in N2 or air to temperatures of 200–300  C. While short-bridged precursors such as BTESM seem to have a higher thermal stability up to 400  C owing to the difficulty for oxidant (O2) to enter the matrix and for reaction products such as CO, CO2, etc. to leave the micropores, higher temperatures may lead to partial oxidation or degradation of the organic bridge and are generally avoided. The macroporous α-alumina support on which the hybrid film is deposited may be a disk or a tube. α-Alumina supports have a thickness of about 1–2 mm, a porosity of 30–35%, and an average pore size of the order of 100 nm. The pore size provides a rough estimate of the surface roughness. Since the final hybrid silica separation layer has a desired thickness of only 50–300 nm, the coarse α-alumina surface needs to be smoothened before the silica film can be applied. An intermediate γ-alumina layer of 1–3 μm thickness is therefore deposited on the α-alumina support. The layer is prepared from a boehmite (AlOOH) sol by a dipcoating process. The γ-alumina phase combines high porosity (55%) with a narrow mesoporous pore size distribution with a pore size of 3–6 nm. The film is thermally treated at 400–600  C in air and subsequently used as a substrate for hybrid silica membrane deposition. Figure 4 shows a BTESE-based membrane on a mesoporous γ-alumina layer that is coated on a coarse-grained macroporous α-alumina support structure.

Fig. 4 Cross-sectional scanning electron micrograph of a hybrid composite membrane, showing the support, mesoporous cushion layer, and the selective hybrid silica top layer which has a thickness of 100 nm

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Determination of Membrane Pore Size Distribution The most direct way to estimate or determine the pore size distribution of a microporous membrane is by performing single gas permeation measurements with a series of spherical probe molecules of known kinetic diameter. Ideally, the probes should be inert, i.e., have no or very weak (van der Waals) interactions with the pore wall, so that differences in permeability between different probes are solely determined by differences in their size. The permeability of these probe molecules as a function of their kinetic diameter provides a good estimate of the real pore size distribution of the membrane. Of course, many small molecules do not show ideal gas behavior. Simple gases such as He (0.20 nm) and hydrogen (0.29 nm) are relatively inert at room temperature and above, but gases such as CO2 (0.33 nm) can adsorb on the inner pore walls and even show surface diffusion rather than gas translation. Gases and vapors like CH4 (0.38 nm), C2H6 (0.42 nm), i-C4H10 (0.54 nm), and SF6 (0.55 nm) are commonly used (Kanezashi et al. 2014). The permeances of a BTESE, a BTESM, and a silica membrane as a function of kinetic diameter shown in Fig. 5 give a semiquantitative description of the (differences in) pore size distribution between the three systems. The hybrid membranes have a higher permeability than silica at any given kinetic diameter. The slope of the curves is a measure of the selectivity of the membranes for separation of mixtures of varying molecular sizes. The steeper slope of curve of the silica membrane indicates its higher permselectivity for small versus larger gases. For pore size determination of mesoporous (>2 nm pores) membranes, the so-called permporometry technique is commonly used. Permporometry yields the pore size distribution of a (supported) membrane and provides an estimate of Fig. 5 Gas permeance versus molecular kinetic diameter of gases for BTESE, BTESM, and inorganic silica (Tsuru et al. 2006) membranes at 200  C (closed symbols, without heat treatment; open symbols, with heat treatment at 400  C in air for 3 h) (Reprinted with permission from Kanezashi et al. (2012a). Copyright 2012 American Chemical Society)

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the concentration of defects in the layer. Essentially, the method is based on the use of the Kelvin equation for capillary condensation, ln( p/p0) = 2γsVm/rKRT, where p/p0 is the relative pressure of some condensable vapor (e.g., cyclohexane), γs the surface tension of the condensed phase in the pore, rK the so-called Kelvin radius, Vm the molar volume of the condensed phase, R the gas constant, and T the temperature (in K). When a (meso)porous material is exposed to cyclohexane vapor with a set pressure p/p0, all pores in the matrix with radius < rK will be filled by liquid cyclohexane via capillary condensation. The gas flux of another, inert gas, will only occur through those pores that are still open, i.e., those that have a pore radius > rK for a given p/p0. Hence, the magnitude of the flux is a measure of the fraction of pores with radius > rK. By systematically varying p/p0 of the condensable vapor, a pore size distribution can be extracted from the data. An inert gas flux at p/p0 = 1 is an indication of the presence of large defects in the film. When the flux is zero under such conditions, the membrane is considered defect-free. Unfortunately, the permporometry analysis is only valid under conditions where the Kelvin equation is valid, i.e., when the condensed (cyclohexane, water) phase can be considered as a continuous medium. Obviously, molecules in micropores do not form a continuous phase, and so the method is strictly speaking not applicable to microporous systems. But in practice, the Kelvin equation can still be extrapolated to obtain a rough measure of pore size, in particular for the qualitative comparison between similar systems. A method proposed and applied by Tsuru and coworkers is to determine at which relative pressure of the condensable gas the inert gas flux is only 50% of its value under dry conditions. Using the original Kelvin equation, an “effective Kelvin radius” (rather than a complete pore size distribution) is calculated from the data (Tsuru et al. 2003). Typically compounds like water and helium may be used in these experiments as condensable vapor and inert gas, respectively, because of their small kinetic diameters, which make them suitable for microporous matrices. The disadvantage of using water is that it is a rather polar molecule so that its condensation may be influenced by additional polar interactions not taken into account by the Kelvin equation.

Hybrid Organosilica Membranes for Gas Separation and Pervaporation For a membrane to operate properly in a separation process, it needs to have three main properties: 1. Have a high permeability for the selected species in order to maximize the yield of the process 2. Serve as a selective barrier, excluding one or more components of a gas or liquid mixture, while allowing other selected species to pass 3. Have a long lifetime under the conditions of the separation process, i.e., the chemical and thermal stability should be sufficient, even under cleaning operations that typically make use of high pH chemicals

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The ability of a species i to penetrate a membrane matrix is expressed in terms of the permeance Pi. Its value is proportional to the concentration ci and mobility bi inside the microporous matrix, i.e., Pi  bici. The permeance is essentially the ratio between flux ji of species i and applied driving force across the membrane: ji ¼ Pi

  @ci @x

Here ð@ci =@xÞ refers to the concentration gradient of species i, assuming a one-dimensional flux in x-direction. The membrane selectivity depends on variations in the ability of various species to travel through a membrane matrix. For a feed mixture of two gases with the same partial pressure, the so-called ideal permselectivity Si,j of a membrane is defined as: Si, j ¼ Pi =Pj




The ideal permselectivity provides a rough indication of the separating capacity of a membrane material for a given mixture. Its value would have practical meaning if the respective fluxes of components i and j from a mixture would not be interfering with each other during their transport through the microporous matrix, as they do in reality. In practice, the individual permeances are determined via single-component gas flux measurements, and the ideal permselectivity is calculated from those data. For a molecular sieving membrane, the pore size should be of similar size as the molecular diameter. It was already shown for gas transport in microporous TEOSderived membranes that the most important parameter determining the permeance of a species is the membrane pore size (de Vos et al. 1999; Nair et al. 2000). However, since the pore wall is very near to the molecule at all times, the inner pore chemistry also affects the transport of some or all passing species (Boffa et al. 2008; Sekulic et al. 2002). The membrane permselectivity is determined by the width of the pore size distribution; the narrower the pore size distribution, the more selective the membrane can be. The chemical, thermal, and mechanical stabilities of a membrane are determined by both chemical composition and the 3D morphology of the microporous matrix. To achieve true microporosity, i.e., pore sizes 99.5 wt% water in the permeate stream (αH2O/BuOH > 2000) in later work (Kreiter et al. 2009). The selectivity was explained by the very small pore size of BTESE–MTES, as discussed in the previous section. The first purely BTESE-based hybrid silica membrane has considerably larger pores (C2H2 adsorption 514–546 m2/g; N2 131–311 m2/g) than the aforementioned compositions (Castricum et al. 2008a), and such a membrane would therefore be expected to be less selective. However, the separation factor in water/n-butanol pervaporation (αH2O/BuOH = 360–2700) was higher than that of BTESE–MTES, showing that besides the pore size, also other factors are important, e.g., pore connectivity and physicochemical interactions between permeating species and membrane matrix. BTESE membranes are resistant to aggressive aprotic solvents like N-methyl pyrrolidone (NMP), organic acids (Tsuru et al. 2012; van Veen

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Fig. 6 Long-term separation (stars, water content of permeate; squares, water flux; diamonds, nbutanol flux) of a BTESE membrane from 5 wt% water/95 wt% n-butanol binary liquid at various HNO3 concentrations at 95  C. Arrows indicate moments at which additional HNO3 was added to keep the concentration constant (Reprinted from Castricum et al. (2008a) with permission from Elsevier)

et al. 2011), and HNO3 up to concentrations of 0.05 wt% in water–ethanol (pH 2.2) (van Veen et al. 2011). Higher concentrations of HNO3, or the use of stronger acids like methyl sulfonic acid, lead to fast deterioration of the membrane and complete loss of performance within a number of days (Castricum et al. 2008a; van Veen et al. 2011). See Fig. 6. The decreasing water flux and increasing butanol fluxes indicate that water-selective small micropores are disappearing at very low pH, while the contribution of nonselective wider pores that are accessible to n-butanol is increasing. Even higher water fluxes and higher separation factors were obtained when the -C2H4- bridge (BTESE) was replaced by -CH2- (BTESM) (Kreiter et al. 2009, 2011). The pore size of BTESM-derived membranes is smaller than BTESE but larger than BTESE–MTES (Kreiter et al. 2009). More importantly, the CHx:Si ratio of BTESM is only 0.5, whereas that of BTESE and BTESE–MTES is 1. This renders the BTESM membrane a more hydrophilic character than BTESE, MTES, or BTESE–MTES, and the high water flux may at least be partly attributed to that hydrophilic nature. The BTESM membrane also showed a surprising ability to separate water from lower alcohols, including the very difficult molecular separation of methanol and water (Kreiter et al. 2009; ten Elshof et al. 2003). A 5:95 wt/wt water–methanol mixture at the feed side of a BTESM membrane yielded a permeate stream containing 55:45 wt/wt water–methanol. This corresponds to a separation factor αH2O,MeOH = 23. This may seem low compared to values >1000 reported for water/n-butanol (Gallego-Lizon et al. 2002; Sekulic et al. 2002), but methanol and

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water are very similar in size and comparable in their polarity and hydrogenbonding capability (ten Elshof et al. 2003). No other ceramic membranes are known that can separate water from methanol effectively while maintaining a high flux. For example, BTESE–MTES and BTESE are not selective for methanol–water, and even for ethanol–water, the BTESE–MTES membrane only has low selectivity (αH2O,EtOH = 15) (Kreiter et al. 2011). BTESM has separation factors >150 for ethanol, i-propanol, and n-butanol and shows higher water fluxes under otherwise similar conditions. The degree of hydrophobicity of the membrane matrix can be controlled by adjusting the CHx:Si ratio. Obviously, longer aliphatic or aromatic bridges between the silicon centers may lead to membranes with a larger average pore size. Moreover, the network may become more flexible owing to the flexibility of longer chains, in particular long-chain alkylenes. A pervaporation study using hybrid silica membranes for water/n-butanol made from bridged silsesquioxane precursors with alkylene -CnH2n- (n = 1, 2, 8) and arylene (-p-C6H4- and -p-C6H4-p-C6H4-) bridges clearly demonstrated that the n-butanol flux goes up when the CHx:Si ratio is increased from 1 to 6 or 12, while the water flux decreases (Castricum et al. 2011). The phenylene- and biphenylene-bridged systems led to lower increases of the n-butanol permeance than alkylene bridges with similarly high CHx:Si ratios did. The effect of incorporated terminal alkyl groups CnH2nþ1 with n = 1, 2, 3, 6, and 10 was demonstrated in a follow-up study (see Fig. 7). An increasing separation factor for n-butanol was seen when the chain length was increased, with a separation factor αBuOH/H2O = 14 for CHx:Si = 5 (Paradis et al. 2013).

Fig. 7 Permeate concentration of BuOH and H2O as function of the number of C atoms in the terminal R group in pervaporation of n-butanol/water feed mixtures of 95/5 and 5/95 wt%. The feed concentrations were normalized to 5 wt% for direct comparison (Reprinted from Paradis et al. (2013) with permission from Elsevier)

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Gas Separation The obvious difference between gas and liquid separations is the much smaller concentrations in the feed stream in the case of gas separation (Kanezashi et al. 2009). The size differences between molecules in the gas phase are usually smaller than in the liquid phase, while polar (or even electrostatic) interactions between pore wall and gas molecule are absent in many cases. This is the reason that many gas separations work on the principle of molecular sieving by size selection, in particular for industrially relevant nonpolar gases like H2, N2, CH4, and CO (Agirre et al. 2014). The kinetic diameters of small gases differ less than 0.1 nm from each other (H2 0.29 nm, CO2 0.33, N2 0.36, CO 0.37, CH4 0.38 nm), so a very sharp pore size cutoff is required to obtain highly selective membranes. An exception is carbon dioxide, which behaves differently from the others. CO2 is thought to diffuse through the porous network via a surface diffusion mechanism that involves the adsorption of the CO2 molecule (Castricum et al. 2011). Very high separation factors have been reported for amorphous microporous silica, even up to H2/N2 > 4000 (de Vos and Verweij 1998). MTES-derived methylated silica membranes show considerably lower permselectivities (de Vos et al. 1999). Permselectivities have also been quite moderate for gas separation by hybrid organosilica membranes (Kanezashi et al. 2009), and a considerable improvement is required to achieve comparable permselectivities as were already obtained with microporous silica (Kanezashi et al. 2009; Kreiter et al. 2011). The low permselectivity is attributed to the relatively larger pore sizes of -C2H4and -C2H2- bridged silicas compared to microporous SiO2. As discussed above, a small but considerable fraction of pores with a pore diameter > 0.4 nm is present. Since most gases are < 0.4 nm, the theoretically achievable permselectivity for small gases is limited. Even BTESM-derived membranes only have slightly higher selectivity than BTESE-derived membranes in H2 separation (see Fig. 5). Reported ideal permselectivities are 15–21 for H2/N2 and 7–9 for H2/CH4 (Kreiter et al. 2011). BTESM is selective in the separation of hydrocarbons with kinetic diameters > 0.4 nm, e.g., C2H6 (0.42 nm), i-C4H10 (0.54 nm) (Kanezashi et al. 2013), and the separation of propene (0.46 nm) from propane (0.50 nm) (Kanezashi et al. 2012a, b). In agreement with expectation, longer and stiffer bridges lead to larger pores and therefore to lower permselectivities (Castricum et al. 2011; Xu et al. 2014). It was shown recently that the gas selectivity of BTESE membranes can be improved considerably to a permselectivity H2/N2 > 400 by careful engineering of the pore structure during membrane fabrication (Castricum et al. 2015). By employing sols with a low acid ratio ([H+]/[Si] = 0.01) and by coating them onto supports that had been pre-dried at low relative humidity (RH = 0.5%), a hybrid membrane network was formed that contained a considerably lower fraction of pores > 0.35 nm than membranes obtained at higher acid ratios and/or higher humidity (see Fig. 8). The permeance of gases < 0.35 nm like He, H2, and CO2 was not affected by changes in the membrane preparation process. It is hypothesized that both the absence of a high concentration of protons (Castricum et al. 2014) and the absence of water lead to more densified silica network in which larger pores are

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Fig. 8 Normalized single gas permeances of BTESE-based membranes prepared with acid ratio [H+]/[Si] = 0.01 or 0.1, coated onto support systems pretreated at RH = 0.5% (“dry”) or 90% (“moist”). Reference permeance is hydrogen flux at 473 K, Δp = 2 bar (Reprinted from Castricum et al. (2015) with permission from Elsevier)

virtually absent. Formation of larger pores can thus be understood as resulting somehow from a higher condensation rate and/or longer drying times when abundant water is present. Next to hydrogen separation membranes, there is also a clear need for CO2selective membranes, as upcoming CO2 sequestration technologies require energyefficient separation of CO2 from (exhaust) gases. However, well-performing membranes are lacking. Unlike most nonpolar gases, the activation energy of CO2 permeance in silica and methylated silica membranes is typically below zero, and this is also the case in hybrid organosilica (Castricum et al. 2011). The activation energy of permeance is essentially the activation energy of diffusion minus the heat of sorption of the gas molecule on the pore wall (ten Elshof et al. 2003). For nonpolar gases, the temperature dependency of permeance is usually positive or slightly negative, i.e., values between 1 and +2 kJ/mol are normal. The more negative activation energy for CO2 suggests that its transport path is influenced by chemical interactions (exothermic adsorption) with the hybrid silica matrix. Since the activation energy of hydrogen permeance is positive, a CO2 separation membrane operates ideally at an as low as possible operating temperature so that adsorption of CO2 is maximal, while H2 permeance is minimal. For most other separations, higher operating temperatures are usually desired because of higher fluxes and improved selectivities. The permeance of CO2 also decreases with increasing gas pressure, due

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to partial pore blocking resulting from pore filling with condensed CO2 (Castricum et al. 2011). The permeances of the other, inert gases are not influenced by pressure. Dispersion of niobium (Nb) into a microporous silica membrane (Nb:Si = 1:3) has been shown to suppress the transport rate of CO2 (Boffa et al. 2008). This was explained by the formation of acidic Nb sites in the matrix structure leading to strong gas adsorption. The same approach was also pursued for BTESE-derived silica (Qi et al. 2010, 2012; Qureshi et al. 2015), but in these cases, the phenomenon was either observed or not, depending on the pretreatment of the BTESE membrane. Low CO2 permeances were observed when the system was annealed at high temperatures (400–550  C in N2) (Qi et al. 2012) but not when lower annealing temperatures (300  C in N2) were employed (Qureshi et al. 2015). BTESE is known to be stable at 300  C, but (partial) thermal degradation of the organic bridges may occur above 400  C (Agirre et al. 2014; Qi et al. 2012), and this may have influenced the permselectivity of the membrane toward CO2 in those respective cases. Other dopants in BTESE and BTESM matrices that have been tested are Al (Kanezashi et al. 2013), B, and Ta (all 300  C in N2) (Qureshi et al. 2015), but also in these cases, no specific influence on CO2 permeability was observed. However, both transition metal doping and the use of low dip sol concentrations during membrane fabrication contribute to the formation of more gas-permeable membranes in general (Qureshi et al. 2013, 2015). In general, the same general trends in gas permeation behavior are observed in hybrid organosilicas with varying alkylene and phenylene bridges as in silica and methylated silica (MTES derived), irrespective of the type of gas and the type of bridging group. The only exception is the -C8H16- bridged membrane, which shows relatively high activation energies of gas permeance (+6 kJ/mol for He, H2, N2, and CH4 and +2 kJ/mol for CO2) (Castricum et al. 2011). Since the end-to-end lengths in a biphenylene-bridged precursor are similar to that of an octylene bridge, the difference is likely associated with the stiffness of the bridging group. The long and flexible -C8H16- bridges seem to have a retarding effect on the transport rate of all gases, probably as a result of pore blocking. The effect increases with temperature, probably due to thermal vibrations and conformational changes of the octylene bridge that slow down gas transport further. Shorter and more rigid bridges are apparently beneficial for high transport rates. Relatively high permselectivities were observed for C8H16-, phenylene-, and biphenylene-bridged membranes. The larger pore sizes seem to be beneficial for affinity-based selection at high temperature (Castricum et al. 2011). Hence, the permeability of gases depends on the size, stiffness, and nature of the organic bridging group, and in this way, the selectivity of hybrid silica membranes can be tailored toward certain targeted molecules. Recently, more complex functional groups to tailor the transport properties of hybrid membranes have been reported. A malonamide-bridged hybrid membrane with an ability to disperse transition metal ions homogeneously throughout the matrix has been reported (Besselink et al. 2015). As examples both Ce4+ and Ni2+ were doped into a hybrid organosilica membrane. Also a triazine-functional hybrid silica membrane based on a novel solgel precursor with a high H2/SF6 selectivity and

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a high affinity for propene/propane separation has been reported (Ibrahim et al. 2014a, b). Functionalization of the hybrid matrix by co-condensation of BTESE with an amine-functional silane precursor has also been reported (Paradis et al. 2012).

Conclusions Hybrid organosilica membranes based on solgel-processed-bridged silsesquioxane precursors are the first generation of oxide-based molecular separation membranes that show great promise for actual application in industrially relevant gas and liquid separation processes. Since the number of possible bridging groups is virtually infinite, many other functional groups may be covalently incorporated into the hybrid silica matrix, resulting in novel molecular sieving membranes with yet unexplored and possibly unprecedented performance in certain separations. We expect to see new types of membranes based on this principle in the forthcoming years.

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Kanezashi M, Shazwani WN, Yoshioka T, Tsuru T. Separation of propylene/propane binary mixtures by bis(triethoxysilyl) methane (BTESM)-derived silica membranes fabricated at different calcination temperatures. J Membr Sci. 2012b;415:478–85. Kanezashi M, Miyauchi S, Nagasawa H, Yoshioka T, Tsuru T. Pore size control of Al-doping into bis (triethoxysilyl) methane (BTESM)-derived membranes for improved gas permeation properties. RSC Adv. 2013;3(30):12080–3. Kanezashi M, Miyauchi S, Nagasawa H, Yoshioka T, Tsuru T. Gas permeation properties through Al-doped organosilica membranes with controlled network size. J Membr Sci. 2014;466: 246–52. Keizer K, Uhlhorn RJR, Vanvuren RJ, Burggraaf AJ. Gas separation mechanisms in microporous modified gamma-Al2O3 membranes. J Membr Sci. 1988;39(3):285–300. Kreiter R, Rietkerk MDA, Castricum HL, van Veen HM, ten Elshof JE, Vente JF. Stable hybrid silica nanosieve membranes for the dehydration of lower alcohols. ChemSusChem. 2009;2 (2):158–60. Kreiter R, Rietkerk MDA, Castricum HL, van Veen HM, ten Elshof JE, Vente JF. Evaluation of hybrid silica sols for stable microporous membranes using high-throughput screening. J Sol-Gel Sci Technol. 2011;57(3):245–52. Leenaars AFM, Keizer K, Burggraaf AJ. The preparation and characterization of alumina membranes with ultra-fine pores. J Mater Sci. 1984;19(4):1077–88. Lin MY, Lindsay HM, Weitz DA, Ball RC, Klein R, Meakin P. Universality in colloid aggregation. Nature. 1989;339(6223):360–2. Lin YS, Kumakiri I, Nair BN, Alsyouri H. Microporous inorganic membranes. Sep Purif Methods. 2002;31(2):229–379. Loy DA, Carpenter JP, Myers SA, Assink RA, Small JH, Greaves J, et al. Intramolecular condensation reactions of alpha, omega-bis(triethoxysilyl)alkanes. Formation of cyclic disilsesquioxanes. J Am Chem Soc. 1996;118(35):8501–2. Lu GQ, da Costa JCD, Duke M, Giessler S, Socolow R, Williams RH, et al. Inorganic membranes for hydrogen production and purification: a critical review and perspective. J Colloid Interface Sci. 2007;314(2):589–603. Maene N, Nair BN, D’Hooghe P, Nakao SI, Keizer K. Silica-polymers for processing gas separation membranes: high temperature growth of fractal structure. J Sol-Gel Sci Technol. 1998;12(2):117–34. Nair BN, Keizer K, Suematsu H, Suma Y, Kaneko N, Ono S, et al. Synthesis of gas and vapor molecular sieving silica membranes and analysis of pore size and connectivity. Langmuir. 2000;16(10):4558–62. Ockwig NW, Nenoff TM. Membranes for hydrogen separation. Chem Rev. 2007;107(10): 4078–110. Paradis GG, Kreiter R, van Tuel MMA, Nijmeijer A, Vente JF. Amino-functionalized microporous hybrid silica membranes. J Mater Chem. 2012;22(15):7258–64. Paradis GG, Shanahan DP, Kreiter R, van Veen HM, Castricum HL, Nijmeijer A, et al. From hydrophilic to hydrophobic HybSi (R) membranes: a change of affinity and applicability. J Membr Sci. 2013;428:157–62. Prabhu AK, Oyama ST. Highly hydrogen selective ceramic membranes: application to the transformation of greenhouse gases. J Membr Sci. 2000;176(2):233–48. Qi H, Han J, Xu NP, Bouwmeester HJM. Hybrid organic-inorganic microporous membranes with high hydrothermal stability for the separation of carbon dioxide. ChemSusChem. 2010;3 (12):1375–8. Qi H, Chen HR, Li L, Zhu GZ, Xu NP. Effect of Nb content on hydrothermal stability of a novel ethylene-bridged silsesquioxane molecular sieving membrane for H2/CO2 separation. J Membr Sci. 2012;421:190–200. Qureshi HF, Nijmeijer A, Winnubst L. Influence of sol-gel process parameters on the microstructure and performance of hybrid silica membranes. J Membr Sci. 2013;446:19–25. Qureshi HF, Besselink R, ten Elshof JE, Nijmeijer A, Winnubst L. Doped microporous hybrid silica membranes for gas separation. J Sol-Gel Sci Technol. 2015;75(1):180–8.

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Sekulic J, Luiten MWJ, ten Elshof JE, Benes NE, Keizer K. Microporous silica and doped silica membrane for alcohol dehydration by pervaporation. Desalination. 2002;148(1–3):19–23. Sekulic J, ten Elshof JE, Blank DHA. A microporous titania membrane for nanofiltration and pervaporation. Adv Mater. 2004;16(17):1546–50. Shea KJ, Loy DA. A mechanistic investigation of gelation. The sol-gel polymerization of precursors to bridged polysilsesquioxanes. Acc Chem Res 2001;34(9):707–16. Shimoyama T, Yoshioka T, Nagasawa H, Kanezashi M, Tsuru T. Molecular dynamics simulation study on characterization of bis(triethoxysilyl)-ethane and bis(triethoxysilyl)ethylene derived silica-based membranes. Desalin Water Treat. 2013;51(25–27):5248–53. Spijksma GI, Huiskes C, Benes NE, Kruidhof H, Blank DHA, Kessler VG, et al. Microporous zirconia-titania composite membranes derived from diethanolamine-modified precursors. Adv Mater. 2006;18(16):2165–8. ten Elshof JE, Abadal CR, Sekulic J, Chowdhury SR, Blank DHA. Transport mechanisms of water and organic solvents through microporous silica in the pervaporation of binary liquids. Microporous Mesoporous Mater. 2003;65(2–3):197–208. Tsuru T, Takata Y, Kondo H, Hirano F, Yoshioka T, Asaeda M. Characterization of sol-gel derived membranes and zeolite membranes by nanopermporometry. Sep Purif Technol. 2003;32(1–3): 23–7. Tsuru T, Shintani H, Yoshioka T, Asaeda M. A bimodal catalytic membrane having a hydrogenpermselective silica layer on a bimodal catalytic support: preparation and application to the steam reforming of methane. Appl Catal A Gen. 2006;302(1):78–85. Tsuru T, Shibata T, Wang JH, Lee HR, Kanezashi M, Yoshioka T. Pervaporation of acetic acid aqueous solutions by organosilica membranes. J Membr Sci. 2012;421:25–31. Uhlhorn RJR, Keizer K, Burggraaf AJ. Gas and surface diffusion in modified gamma-alumina systems. J Membr Sci. 1989;46(2–3):225–41. Uhlhorn RJR, Keizer K, Burggraaf AJ. Gas-transport and separation with ceramic membranes. 2. Synthesis and separation properties of microporous membranes. J Membr Sci. 1992;66 (2–3):271–87. van Veen HM, Rietkerk MDA, Shanahan DP, van Tuel MMA, Kreiter R, Castricum HL, et al. Pushing membrane stability boundaries with HybSi (R) pervaporation membranes. J Membr Sci. 2011;380(1–2):124–31. Wei Q, Wang F, Nie ZR, Song CL, Wang YL, Li QY. Highly hydrothermally stable microporous silica membranes for hydrogen separation. J Phys Chem B. 2008;112(31):9354–9. Xu R, Ibrahim SM, Kanezashi M, Yoshioka T, Ito K, Ohshita J, et al. New insights into the microstructure-separation properties of organosilica membranes with ethane, ethylene, and acetylene bridges. ACS Appl Mater Interfaces. 2014;6(12):9357–64. Zhang Y, Sunarso J, Liu SM, Wang R. Current status and development of membranes for CO2/CH4 separation: a review. Int J Greenhouse Gas Control. 2013;12:84–107.

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Shridevi Shenoi-Perdoor, Achraf Noureddine, Fabien Dubois, Michel Wong Chi Man, and Xavier Cattoën

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Click Reactions for Sol-Gel Precursors and Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thiol-ene Click Reactions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . The CuAAC Reaction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Miscellaneous Click Reactions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . CuAAC Versus Thiol-ene Chemistry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Combination of Various Click Reactions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Technological Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Application in Separation Science . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Application in Catalysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Application in Biology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

The design of complex, multifunctional, and versatile materials for diverse applications ranging from nanomedicine to catalysis requires simple and reliable functionalization methodologies. Click chemistry describes a family of efficient, simple, and wide-scope reactions that can easily be applied to sol-gel materials. This chapter focuses on the use of click chemistry on various silica-based materials. Thiol-ene and copper-catalyzed alkyne to azide cycloaddition S. Shenoi-Perdoor · F. Dubois · X. Cattoën (*) University of Grenoble Alpes, Inst NEEL, Grenoble, France CNRS, Institut NEEL, Grenoble, France e-mail: [email protected] A. Noureddine · M. Wong Chi Man Institut Charles Gerhardt Montpellier, UMR-5253, CNRS, ENSCM, Université Montpellier, Montpellier, France # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_95

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(CuAAC), the two most popular examples of click reactions, are described extensively, with prime focus on their use for the functionalization of sol-gel materials such as bulk mesoporous silica, silica monoliths, mesoporous silica nanoparticles (MSNs), and periodic mesoporous organosilicas (PMOs). In addition to the mechanism of functionalization, reaction conditions are highlighted. Furthermore, the applications of these click-functionalized materials in diverse fields are discussed to highlight the vast potential and the numerous advances rendered possible thanks to these methodologies.

Introduction The controlled functionalization of sol-gel materials with organic functionalities is of paramount importance for the development of functional materials for various applications such as optics, coatings, sensing, medical devices, or catalysis (Sanchez et al. 2011; Sanchez 2005). For each of these applications, one must select an optimized inorganic support together with the targeted organic functionality and must be able to connect the two entities through inert chemical bonds. While sol-gel technologies have been continuously progressing to enable the preparation of metal-oxide materials under various shapes (bulk powders, monoliths, fibers, nanoparticles, thin films, etc.), textures, and porosities (notably with the advent of ordered mesoporous materials) (Yanagisawa et al. 1990; Kresge et al. 1992), the incorporation of organic moieties within the materials still suffered from strong limitations. Indeed, the covalent immobilization was limited to the following methods, illustrated here for silica-based materials (Fig. 1): (i) the grafting of organo(alkoxy)silanes or organo(chloro)silanes on preformed materials, (ii) the sol-gel co-condensation of a silica precursor with an organo(alkoxy)silane, and (iii) the formation of pure organo(silsesquioxanes) and in particular bridged silsesquioxanes (Hoffmann et al. 2006; Kickelbick 2004). Furthermore, some studies relied on the use of chemical reactions on reactive functionalities previously incorporated within the material through one of the previous methods, but most often under rather harsh conditions or lacking a good control on the conversion or on the selectivity. These limitations are now being overcome by the recent advent of the so-called click reactions. Click chemistry, as first described by Sharpless in 2001 (Kolb et al. 2001), encompasses all reactions that enable easy linking of two reactive organic entities provided they are of wide scope, modular, and selective and occur with high yields and under mild and green conditions using harmless reactants and benign solvents. This can in fact be related to Lego chemistry, where a simple reaction enables easy covalent linking of any kind of organic moiety. While this concept was initially established in the context of synthetic organic chemistry for drug discovery, it was quickly extended and applied to polymer science and as a reliable functionalization method in materials chemistry, in particular for sol-gel materials (Cattoën et al. 2014).

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Fig. 1 Common routes and click grafting to functional organosilicas (Adapted from Ferré et al. 2016 with permission from The Royal Society of Chemistry)

The advantages of applying click reactions to sol-gel materials are numerous (section “Click Reactions for Sol-Gel Precursors and Materials”). As a postfunctionalization method, click chemistry allows to incorporate functional moieties while preserving the inorganic skeleton of the sol-gel material, thus the textural and morphological properties; the wide scope and the mild conditions associated with the click reactions enable the covalent incorporation of biomolecules such as enzymes or DNA fragments hardly grafted by the previous methods; the incorporation of the clickable moieties by co-condensation allows to randomly position the functions throughout the material, thus a precise design of the materials. Thanks to all these advantages, click reactions are now widely applied for the preparation of functional materials to be used in supported catalysis, sensing, separation, or nanomedicine (section “Technological Applications”).

Click Reactions for Sol-Gel Precursors and Materials The most common examples of click chemistry are the Cu(I)-catalyzed alkyne–azide cycloaddition to form a 1,2,3 triazole linkage, the thiol-ene reaction between a thiol and an alkene, and to a less extent reactions such as epoxide ring openings or oxime ligation.

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Fig. 2 Radical (a) and ionic (b) thiol-ene click reaction mechanisms

Thiol-ene Click Reactions The thiol-ene coupling involves the addition of a thiol function to a terminal double bond, which can be activated or not by the presence of an electron-withdrawing group (EWG). Terminal double bonds and thiols are relatively easy to produce and are present in a myriad of natural or synthetic functional molecules, which enables a wide use for materials functionalization.

Mechanisms of the Thiol-ene Couplings Depending on the reaction conditions, the addition of a thiol to an alkene can follow two different mechanisms, either radical or ionic (Lowe 2010). Most thiol-ene reactions proceed under radical pathways (Fig. 2a), via thiyl radicals generated either under UV irradiation, or using an initiator. The thiol-ene reactions proceeding via a nucleophilic mechanism (Fig. 2b) are often conducted in the presence of a base and using activated olefinic substrates (acrylates or maleimides). These reactions may be catalyzed by strong nucleophiles such as phosphines or amines. Interestingly, both mechanisms feature an anti-Markovnikov regioselectivity leading to linear adducts. Sol-Gel Precursors by Thiol-ene Coupling The most popular alkoxysilane derivatives enabling functionalization by thiol-ene coupling include mercaptopropyltrialkoxysilanes (MPTMS, MPTES) and allyl- or vinyltrialkoxysilanes (ATES, VTES) as well as γ-methacryloxypropyltriethoxysilane (γ-MAPS) for the complementary retrosynthesis (Fig. 3a). Furthermore, a number of diverse (organo)trialkoxysilanes with targeted functionalities can be synthesized using thiol-ene click chemistry in nearly quantitative yields with high purity. This has been done either by reacting alkenes with MPTES or MPTMS or by reacting thiols with allyltrialkoxysilanes, using Irgacure 651 as the photoinitiator under UV irradiation (Fig. 3b) (Tucker-Schwartz et al. 2011). This led to different interesting sol-gel precursors featuring diverse functionalities such as thiazole, thiourea, N-Boc-(4R)-4-hydroxy-L-proline methyl ester, quinuclidine, quinine, and guanidinium. Interestingly, a mannose-terminated silane featuring four free hydroxyl groups has been synthesized accordingly (Wu et al. 2013). It is also possible to synthesize alkoxysilanes containing carboxylic acid functions, which are very easily hydrolyzed compounds (Bordoni et al. 2015).

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a

O

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Fig. 3 Common commercial sol-gel precursors (a), typical thiol-ene (b), and thiol-yne click reactions (c)

Photoinitiated radical-based thiol-yne click reactions have also been used to synthesize diverse bridged organosilanes with pendent functional groups (Fig. 3c) (Zuo et al. 2014). This synthesis is performed by reacting alkynes with MPTMS under 100 W UV irradiation in nearly quantitative yields with high purity.

Materials Clickable by the Thiol-ene Coupling Thiol-ene on Bulk Mesoporous Silica The first general study on the use of the thiol-ene reaction to functionalize mesoporous silicas was reported in 2011 by the group of Sen Gupta (Kumari et al. 2011). SBA-15 materials containing activated methacrylate groups introduced by post-grafting were shown to be easily functionalized with a number of thiols using phosphines (TCEP and DMPP) as nucleophilic catalysts, following an ionic mechanism. Interestingly, the reaction can proceed either in aqueous conditions (PBS buffer) or in organic solvents, which allows the click grafting of polar (cysteine, 2-mercaptoethanol, 2-mercaptopropionic acid) as well as nonpolar thiols (6-(Ferrocenyl)hexanethiol) (Fig. 4). The extent of the reaction can be followed by Fourier Transform InfraRed spectroscopy (FTIR) with the shift of the carbonyl stretching vibration frequency of the methacrylate group, while the yields can be deduced by TGA and ICP analyses. Conversions of 45–55% of the grafted acrylate groups were reported. Thiol-ene on Silica Monoliths Clickable monoliths bearing thiol or alkene groups were synthesized by co-condensation between TMOS and a thiol- or alkene-bearing alkoxysilane. On

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Fig. 4 Synthesis of alkene-bearing SBA-15 with the different clicked thiol-bearing compounds (Kumari et al. 2011)

the one hand, monoliths were obtained under acetic acid catalysis using PEG as template (Chen et al. 2011a; Wang et al. 2012a, 2015). On the other hand, monoliths could be obtained without template under basic catalysis, with up to 25 mol% of organosilane in the starting mixture (Göbel et al. 2014). In all cases, most reactive functionalities could be coupled with the corresponding partner under one of the following conditions: (i) without initiator at 60  C for 5 h, (ii) using the V50 initiator at 55  C, and (iii) under 366 nm UV irradiation at 20  C. The occurrence of the functionalization process was usually checked by fluorescence measurements after reaction of the monolith with cysteamine followed by coupling with fluorescein isothiocyanate (FITC). Furthermore, iodine titration of the unreacted double bonds on vinyl or allyl- functionalized monoliths was used to determine conversions in the 25–50% range (Fig. 5a). Alternatively, the click functionalization reaction can be performed simultaneously with the sol-gel preparation of the monolithic material. This allowed for the preparation of glutathione- and acrylamide-containing chromatography columns (Chen et al. 2013; Lin et al. 2014). Thanks to the mild reaction conditions employed, biological species such as proteins and enzymes can be introduced into the columns without deterioration of their activity. For example, the trypsin enzyme was grafted into a γ-MAPS-TMOS hybrid monolithic column by thiol-ene reaction between the monolith and trypsin previously reduced using TCEP (Chen et al. 2011a) (Fig. 5b). Aptamer-based organic-silica hybrid monoliths have also been prepared using a one-pot process in a fused-silica capillary (Wang et al. 2015). Thiol-ene on Silica Nanoparticles Thiol-ene chemistry is also being used for the functionalization of silica NPs for various biological applications. The thiol or alkene functionalities required for the thiol-ene reactions can either be incorporated by co-condensation or by means of grafting on preformed silica particles, which can be easily synthesized by reverse microemulsion or Stöber syntheses. Thereby, several core–shell silica@polymer NPs have been developed, by concomitantly anchoring thiol-terminated polymers (poly(N-isopropylacrylamide) (PNIPam) and poly(2-(diethylamino)ethyl methacrylate) (PDEAMA)) onto silica NPs prepared using Stöber synthesis and surface modified with methacrylate groups using a thermal radical initiator (Kotsuchibashi

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Fig. 5 Preparation of functional monoliths by a post-functionalization (a) or in situ click reaction (b)

Fig. 6 Preparation of thioland allyl-bearing silica nanoparticles and their derivatization with different polymers (Adapted from Kotsuchibashi et al. 2012 with permission from The Royal Society of Chemistry)

et al. 2012) (Fig. 6). The polymer densities at the surface of the NPs were estimated to be between 0.2 and 0.4 nm 2. Fluorescent bioprobes were obtained by anchoring 4-mercaptophenylboronic acid onto vinyl-functionalized silica NPs doped with fluorescein using benzophenone as photoinitiator (Cheng et al. 2013).

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Interestingly, mercapto-functionalized mesoporous silica nanoparticles can be obtained by co-condensation between TEOS and MPTMS and then easily derivatized with maleic anhydride, which allows further transformation into carboxylic acids after saponification. The click grafting procedure involves here a basic mechanism, with the use of trimethylamine as base to promote the thiol addition (He et al. 2014). In addition to MSNs, thiol-ene chemistry can also be applied to periodic mesoporous organosilica (PMO) NPs, which feature the highest possible loading of functionalities (Croissant et al. 2015a). Bifunctional PMO materials bearing sulfonic acid and thiol functionalities were synthesized by the co-condensation of –SH and –SO3H containing phenylene-based bridged organosilanes. The thiol groups on the surface were then used for immobilization of these particles on glass slides treated with alkenes using photoinitiators. The resulting thiol groups enabled further functionalization with Ag nanoparticles (Gehring et al. 2015).

The CuAAC Reaction Among all possible click reactions, the cycloaddition family seems very promising as it allows for the formation of two bonds simultaneously with a cyclic structure. Among them, the Huisgen cycloaddition of azides and alkynes is noteworthy. However, these reactions typically require elevated temperatures and may yield a mixture of regioisomers. In 2002, Meldal and Sharpless independently reported that copper (I) can efficiently catalyze the Huisgen cycloaddition (Meldal and Tornøe 2008; Tornøe et al. 2002; Rostovtsev et al. 2002), yielding a single regioisomer under very mild conditions. The so-called CuAAC reaction underwent intensive research and became extremely popular in various areas of chemistry, expending from drug discovery to polymers and materials chemistries, thanks to the mild and tunable reaction conditions, and the almost infinite substrate scope of this reaction.

Mechanism of the CuAAC Reaction Though a plausible mechanism featuring the formation of a simple copper acetylide and its reaction with organic azides was quickly widely accepted, the mechanism of the CuAAC was only quite recently clarified by V. Fokin and G. Bertrand (Worrell et al. 2013; Jin et al. 2015). Indeed, these studies evidenced the simultaneous participation of two copper centers during the catalytic cycle, as depicted in Fig. 7. One of the primary advantages of CuAAC is that it can be carried out under ambient atmospheric conditions, using both protic and aprotic solvents, including water. This reaction is also tolerant to a large number of functional groups, thereby widening its scope. Cu(I), which catalyzes this cycloaddition and thus eliminates the need to perform such reactions at elevated temperatures (reducing the risk of working with azides), can either be added to the reaction mixture directly or generated in situ.

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Fig. 7 Proposed mechanism of the CuAAC click reaction using two copper centers

While the combination of copper sulfate (CuSO4.5H2O) as the copper source and a reducing agent such as sodium ascorbate (which not only reduces Cu(II) to Cu(I) but also reduces any oxygen species, thus decreasing the formation of oxidative by-products) is mostly preferred, other options such as the use of Cu(I) salts directly (CuI, CuOTf.C6H6, [Cu(NCCH3)4][PF6]) are available. However, when working with such salts, it becomes necessary to perform the reaction under oxygen-free conditions to keep the Cu(I) active as Cu(I) tends to get oxidized to Cu(II) under ambient conditions. A number of ligands such as phosphine and/or amine-based systems have also been studied to complement the CuAAC reaction by stabilizing the active Cu(I) species by forming a ligand–Cu(I) complex, thereby enhancing the rate of the click reaction.

Sol-Gel Precursors by CuAAC The commonly used alkoxysilanes for CuAAC are (3-azidopropyl)trialkoxysilanes (AzPTES, AzPTMS) that have recently become commercially available, while the corresponding alkyne-functionalized silanes must be synthesized (derivatives 1, 2, and 3) (Fig. 8a). The CuAAC can also be used for the synthesis of organoalkoxysilanes with targeted functionality. In this case, anhydrous conditions must be employed, for example, using the CuBr(PPh3)3 catalyst in a dry THF:Et3N mixture (B€urglová et al. 2011, 2014; Moitra et al. 2010). This reaction can also be strongly accelerated by microwave irradiation, yielding functional silanes on the gram scale in only 5 min. Functional groups such as epoxides, amino alcohols, diols, ferrocene, nucleic

Fig. 8 Mono- and bis-silylated precursors containing clickable fragments (a) and their use in producing mono- (b) and poly-silylated functional groups (c)

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bases, dyes, porphyrin, or other photosensitizers have thus been derivatized with alkoxysilane functions (Fig. 8) (Mahtab et al. 2011; Croissant et al. 2015b; Mauriello-Jimenez et al. 2015; Wong Chi Man 2014a, b).

Materials for CuAAC CuAAC on Bulk Mesoporous Silica One of the first CuAAC reactions on mesoporous silica involved the covalent immobilization of the trypsin enzyme in large-pore SBA-15 (Fig. 9) (Schlossbauer et al. 2008). Alkyne-modified trypsin was clicked onto azidegrafted SBA-15 using CuAAC. The immobilization was confirmed using IR spectroscopy and the functionalization density was estimated to be 12 wt% of bound enzyme using TGA, and importantly the enzymatic activity could be preserved thanks to the mild reaction conditions (CuSO4/ascorbic acid system in PBS at 4  C). Systematic studies on the CuAAC click grafting of SBA-15 materials were then reported independently shortly after (Nakazawa et al. 2012; Nakazawa and Stack 2008; Malvi et al. 2009). Materials obtained by grafting azidopropyl groups onto preformed SBA-15 were compared to those obtained by the co-condensation method at different AzPTES loadings. Interestingly, the co-condensation approach was shown to yield hybrid silicas with independent, randomly, but not uniformly distributed functions at low AzPTES loading, while those obtained by grafting showed clustered functionalities (Fig. 10). A wide set of alkynes (propargyl alcohol, phenylacetylene and α-D(+)-propargyl mannopyranoside, azido-glycopyranose, ethynylferrocene, 1-ethynylpyrene and (5-ethynyl-2-pyridylmethylmethyl)bis (pyridylmethyl)-amine) were introduced using the standard CuAAC catalytic mixture (Fig. 9). Notably, only three equivalents of alkyne were necessary to yield 85% conversion over 24 h (Malvi et al. 2009). Propargyl-containing clickable large-pore SBA-15 materials were obtained after grafting the mesoporous material with APTES, then reacting the amino groups with pentynoyl chloride. Low molecular weight PMMA and PEG polymers, as well as D-galactose, were then successfully immobilized using the standard catalytic system. Moderate to high loadings of the polymers were introduced without pore blocking. Interestingly, stoichiometric amounts of the azide-functionalized monosaccharides gave a significant loading of 17 wt% (Huang et al. 2010). The sol-gel process also enables to synthesize highly ordered and perpendicularly oriented mesoporous silica films with different functionalities using a combination of EASA (electrochemically assisted self-assembly) and click chemistry, as demonstrated by the work of Walcarius et al. (Fig. 11) (Vilà et al. 2014). Vertically aligned mesoporous silica was synthesized by EASA using TEOS, AzPTMS, CTAB as the template, and NaNO3 as electrolyte, which resulted in the formation of azide-functionalized hexagonally packed and orthogonally oriented mesoporous channels (MCM-41 type). Interestingly, such electrogenerated thin films present a well-defined oriented structure up to 40% of AzPTMS in the starting sol as opposed to a maximum of 10% in the previous

Fig. 9 Azide-bearing mesoporous silica with a set of ethynylated molecules used to be grafted by conventional CuAAC

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Fig. 10 Illustration of the surface functionalization with conventional grafting (a) and co-condensation methods (b) with their corresponding graphical theoretical distribution (insets) (Adapted from Nakazawa et al. 2012 with permission from the American Chemical Society)

C.E.

N TEOS +

CTAB EtOH/H2O(NaNO3)

H3CO N

Si H3CO

OCH3

+

N

N

–

Precursor species

N+ N OH Si

–1.3V (20 s)

W.E.

–

Surfactant OH

Si

Si

AzPTMS

W.E.

W.E.

Fig. 11 Vertically aligned mesoporous silica films obtained using EASA co-condensation of TEOS and AzPTMS (Adapted from Vilà et al. 2014 with permission from Wiley)

examples. The CuAAC reaction was then used to couple the azide groups with a variety of organic molecules such as ethynylferrocene, propargyl alcohol, 3-ethynylthiophene, and 2-ethynylpyridine, as confirmed using IR spectroscopy and cyclic voltammetry. Mesoporous silica microdot arrays have also been developed using a combination of ink-jet printing (IJP), evaporation-induced self-assembly (EISA), and click chemistry (de los Cobos et al. 2012). Patterns consisting of well-resolved microdots arrays (diameter 96 μm spaced out of 40 μm and height 1.5 μm) with wormlike mesoporosity were obtained by ejection of a sol containing 10% of AzPTES in TEOS using the F127 surfactant (Fig. 12). The pores could be functionalized with propargyl alcohol, methyl pent-4-ynoate, ethynylferrocene, and N-propargyl-4amino-1,8-naphthalimide with nearly complete conversion using the classical catalytic system.

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2.0

25

qz (nm–1)

20 15

1.0

10 0.5

5 0

Intensity (arb. units)

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–1.5 –1.0 –0.5 0.0 0.5 1.0 1.5 qy (nm–1)

(1)

F127

(2)

azide

1.3-triazole

Fig. 12 Clickable 5-layer microdots printed on silica wafer (a) with their wormlike structure evidenced by GISAXS (b) and TEM (c) (Adapted from de los Cobos et al. 2012 with permission from the American Chemical Society)

CuAAC on Bridged Silsesquioxanes and PMOs Bridged silsesquioxanes (BS) featuring pendant alkyne groups were obtained by the basic hydrolysis and condensation of compound 2 (Fig. 8) as aggregates of small NPs (ca 30 nm) (B€ urglová et al. 2014; Noureddine et al. 2014). In spite of the NPs being dense, high conversions were observed for a wide range of azide moieties, as observed by FTIR and solid-state NMR. This material was functionalized with various moieties imparting it either a lipophilic, hydrophilic, or both hydrophobic and lipophobic character as evidenced by nitrogen, water, and cyclohexane sorption experiments (Fig. 13). PMOs synthesized from the same precursor in the presence of the SHS surfactant were also synthesized (B€urglová et al. 2014). Very high conversions were observed, though the mesoporous structure partially collapsed during the reaction. While in the previous example all organic moieties featured clickable functions, thereby leading to massive changes in the materials upon CuAAC grafting, clickable mixed PMOs can also be obtained by the co-condensation strategy, using a “spectator” organo-bridged silane acting as the major structural entity, while functionality is imparted by the clickable silane (Fig. 13). Therefore, azide-functionalized ethanebridged PMOs were obtained from bis(trimethoxysilyl)ethane (BTME) and AzPTES using P123 as the template (Gao et al. 2014a), then derivatized with various alkynecontaining moieties to screen a wide range of surface chemistries while keeping the texture constant for different application. CuAAC on Silica Monoliths Highly porous, clickable silica monolith exhibiting a bimodal pore size distribution (macroporous and organized mesoporous) have been reported for CuAAC derivatization (Keppeler and Husing 2011) (Hierarchically Organization in Monolithic

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Fig. 13 BS and PMOs materials made from either pure clickable precursor (a) or a mixture of spectator and clickable organosilanes (b) (Adapted from Noureddine et al. 2014 and Gao et al. 2014a with permissions from the American Chemical Society and Wiley, respectively)

Sol-gel Materials by H€using, N.). Chloroalkyl-functionalized SBA-15 silica monoliths were first prepared by co-condensation of a tetra(hydroxyethoxy)silane and chloro-methyltrimethoxysilane or (3-chloropropyl)-triethoxysilane using P123 as the structuring agent. The chloro substituents were then converted to azido functionalities followed by CuAAC with various alkynes. Though the methylene linker afforded a higher porosity and better pore organization than with the corresponding propyl linker, lower conversions were obtained due to spatial constraints. In the case of the propyl derivative, almost full conversions were observed for the CuAAC reaction using the Cu(NCCH3)4PF6 catalyst in acetonitrile for 7 days at 60  C. Organosilicas monoliths featuring only clickable organic moieties were prepared by acidic hydrolysis of triisopropoxysilyl derivative 4 (Fig. 14), followed by ammonia-catalyzed condensation in a mold without structuring agent (Schachtschneider et al. 2015). The clickable material was derivatized with a variety of alkyne-containing molecules using Cu(NCCH3)4PF6 in DCM for 16 h at room temperature. Interestingly, when only one extremity of the monolith was dipped in the click solution, a gradient of functionalities could be obtained, whereas two

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Fig. 14 Gradual functionalization of a PMO monolith by CuAAC (Adapted from Schachtschneider et al. 2015 with permission from Wiley)

subsequent dippings from both sides yielded materials with two opposite gradients of functionalities. CuAAC on Silica Nanoparticles The CuAAC reaction has been used on dense silica nanospheres in order to prepare inorganic@polymer core–shell nanostructures in a convergent approach. Various synthetic strategies have been followed to anchor different polymers: In the first report, colloidal silica (75–100 nm) was grafted with (3-bromopropyl) trichlorosilane, followed by nucleophilic substitution to introduce the azide group. An alkyne-terminated clickable polyacrylamide was then introduced using the classical CuAAC conditions in DMSO/water (Ranjan and Brittain 2007). Similar strategies can be followed to prepare SiO2@PS and SiO2@PNIPAM nanocomposites (Ranjan and Brittain 2007; Chen et al. 2011b). Azide groups could be directly grafted onto Ludox silica (20 nm) by reaction with AzPTES in water/ethanol. This was used to graft previously prepared alkyne-modified poly-(L)lysine polymers in tris buffer (Kar et al. 2010). In another approach, the direct preparation of alkynyl-containing silica nanoparticles was successfully achieved thanks to the availability of precursor 1 (Fig. 8a), which was co-condensed with TEOS either using the Stöber synthesis or the reverse microemulsion method. The advantage of having alkyne-modified silica resides in the easier preparation of azide versus alkyne-terminated polymers. This synthetic strategy was used to prepare core–shell fluorescent NPs with PEG moieties (Lu et al. 2009; Tissandier et al. 2012). The synthesis of core–shell silica@polymer nanocomposites can also be performed in a stepwise approach, by grafting reactive monomers at the surface of silica nanoparticles and then performing the polymerization reaction, either RAFT or

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ATRP. In a first report, azide–silica NPs functionalized with trithiocarbamate moieties could thus be submitted to two subsequent polymerization reactions with styrene and then methyl acrylate, yielding SiO2@PS@PMA NPs (Ranjan and Brittain 2008). In a second report, the azide–silica was functionalized with a poly (PEG methacrylate) shell with the simultaneous click grafting of (3-bromopropyl) propynoate and polymerization of PEG methacrylate at 80  C. This was made possible because the CuBr/bipyridine complex is able to catalyze both the CuAAC and the ATRP reactions (Li et al. 2012). Silica NPs containing azide groups were functionalized in a competitive fashion by two different alkyne-derivatized functional groups, bearing either a fluorophore or a biotin moiety, enabling avidin recognition (Achatz et al. 2010). Though the relative loadings of functional groups were not quantified, this competitive approach appears very promising for the multiple functionalization of silica materials.

CuAAC on Mesoporous Silica Nanoparticles The first generalization of the usefulness of the CuAAC reaction to derivatize MSNs with various functionalities was reported in 2011 (Moitra et al. 2011). Clickable MSNs bearing 1% or 10% of azide or alkyne clickable groups were produced according to a method previously reported by S Mann, involving the synthesis of silica nuclei from a concentrated basic solution of CTAB and organosilanes, followed by dilution and neutralization (Lebeau and Innocenzi 2011). This method yielded roughly spherical azide-based MSN, whereas nanospheres and nanorods were obtained with the alkyne precursor depending on its loading. The final amount of clickable groups ranged from 0.16 to 1.1 mmol/g. Interestingly, the functionalization with pyrene derivatives under standard conditions showed almost complete conversions, and the presence of individual, well-separated functions with 1% loading of clickable functions (Moitra et al. 2011). A more straightforward synthesis of azide-based MSNs was then reported (Malvi et al. 2012) based on the typical MSN synthesis procedure first described by V Lin (Lai et al. 2003). This led to further derivatization with an iron–biuret complex to form enzyme mimics (Fig. 15a).

Fig. 15 Different examples using CuAAC click reaction on mesoporous silica nanospheres (a) and hollow nanoPMOs (b) (Adapted from Malvi et al. 2012 and Shi et al. 2011 with permissions from The Royal Society of Chemistry and Wiley, respectively)

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Thanks to the mild conditions offered by the CuAAC reactions, it is possible to perform reactions in water and at low temperature. Therefore, this reaction appeared to be well-suited for the construction of complex gatekeepers based on supramolecular assemblies, which are sensitive to high temperatures and the presence of acids, bases, or organic solvents. The first use of the CuAAC reaction to form nanomachines from MSN was reported by Zink and Stoddart (Patel et al. 2008). After an azide-containing stalk was formed at the surface of amino-functionalized MSNs, rhodamine B was loaded in the pores which were blocked by cyclodextrin which formed a supramolecular complex with the stalk. The pore was then gated with an adamantyl group while keeping intact the rest of the structure thanks to the CuAAC reaction (section “Cancer Therapy”). Similarly, folic-acid terminated gatekeepers were investigated (Porta et al. 2013). The linker between the silica surface and the clickable fragment was also tuned to allow cleavage upon irradiation or in the presence of reducing agents or thiols. Therefore, linkers containing ortho-nitrobenzyl esters were constructed on mesoporous silica nanoparticles to allow the photo-activation of the gates (Park et al. 2009). Furthermore, linkers containing disulfide functions and propargyl groups were immobilized on MSNs allowing glutathione or redoxmediated cleavage (Wang et al. 2012b). The immobilized alkyne fragments were then clicked with azidomethylcyclodextrine. The mild reaction conditions and the wide scope of the CuAAC reaction also enabled the gating of the pores with nucleic acids. Duplex DNA was hence grafted onto azide-functionalized MSNs with the pores filled with rhodamine B, enabling temperature or enzyme-mediated release (Chen et al. 2011c). Notably, MSNs were also clicked at the surface of SBA-15 particles previously loaded with free trypsin, so as to prevent the enzyme from escaping, while enabling the substrates and reactant to diffuse (Malvi and Gupta 2012). PMO nanoparticles have recently emerged as an alternative to MSN, with novel properties arising from the high loading of organic fragments (Croissant et al. 2015a). Hollow PMO nanoparticles have thus been synthesized using a CuAAC click attachment. The parent particles were obtained after co-condensation of BTSB and AzPTMS around hematite cores (100 nm) in the presence of CTAB, followed by acidic etching of the iron oxide core. These materials feature both small mesopores (~3 nm) and big inner voids (~90 nm). They were easily functionalized with alkyne-bearing imidazolidinones using the CuI/THF/DIPEA catalytic system for 3 days at 50  C with ca 80% conversion (Shi et al. 2011) (Fig. 15b).

Miscellaneous Click Reactions In addition to olefins, other compounds such as isocyanates and epoxides are also easily available and very reactive toward thiols, thereby opening up the possibility of alternatives to the conventional click reactions. For instance, mesoporous silica monoliths were synthesized and functionalized using a number of addition reactions: (i) coupling of an isocyanate

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Fig. 16 Miscellaneous click reactions on mesoporous silica monoliths (Adapted from Göbel et al. 2016 with permission from Wiley)

to a surface-immobilized thiol, (ii) addition of an epoxide to a surface-immobilized thiol, (iii) cross-metathesis between two olefins, and (iv) Huisgen [2 + 3] cycloaddition of an alkyne-functionalized monolith with an azide (Fig. 16) (Göbel et al. 2016). It was reported that epoxide and isocyanate additions led to high

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O O Si O

NH2 O

O R H 2O

O O Si O

N O

R

Fig. 17 Functionalization of silica-based materials via oxime ligation

degrees of functionalization while the latter two were less effective. While the reasons for this are still being studied, it is speculated that organometallic catalysis is more difficult to control but this is in contrast to previous work on the CuAAC reaction (Keppeler and H€using 2011). The oxime ligation, i.e., the reaction of alkoxyamines with aldehydes or ketones, is another interesting click reaction that has been shown to be easily applicable to MSNs (Fig. 17). Indeed, MSNs grafted with an alkoxyamine function after protection-deprotection steps were able to react with a range of aldehydes and ketones (Ferris et al. 2015). Interestingly, the oxime linkage can also be cleaved upon acidic hydrolysis. This reaction is very promising, in particular owing to the availability of simple alkoxyamine-bearing sol-gel precursors (Martin et al. 2005).

CuAAC Versus Thiol-ene Chemistry Click chemistry has indeed made possible the easy synthesis of multifunctional surfaces for various applications ranging from catalysis to drug delivery. However, the reaction conditions of the different click reactions impart different characteristics and advantages. CuAAC, which is the most commonly used form of click chemistry, is not only extraordinarily wide in scope but also results in high conversions and can be performed at room temperature, in aqueous as well as organic environments with the formation of chemically stable triazole linkers (Moitra et al. 2010). When azide–silicas are used as substrates, the occurrence of the reaction can easily be monitored by FTIR analysis, which results in simple quantification of the conversion (B€urglová et al. 2014; Malvi et al. 2009). However, its major drawback lies in the toxicity of copper residues on living species, limiting their use for biological applications (Hong et al. 2009). Though several washing strategies are available to flush out the unwanted Cu from the final products, even small traces of Cu(II) could be a major issue for certain biological applications due to the possible production of reactive oxygen species from molecular

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Fig. 18 Cyclooctyne-labeled MSNs are a prospective platform for copper-free click reactions for biomedical applications (Adapted from Lee et al. 2013 with permission from Wiley)

oxygen (Liu et al. 2006). A solution to this problem was proposed by Finn et al. who used a Cu-complexing ligand, tris-(3-hydroxypropyltriazolylmethyl)amine (THPTA), which not only accelerates the CuAAC but also transforms the free copper species into a low-toxicity complex (Hong et al. 2010). Another possible alternative is the use of strained alkynes, namely, cyclooctynes, to allow strainpromoted cycloaddition of alkynes and azides (Baskin et al. 2007). This copperfree Huisgen cycloaddition occurs at room temperature with comparable rates to the CuAAC, but has only been scarcely applied to sol-gel materials (Lee et al. 2013). Indeed, this method involves elaborated cyclooctyne substrates that require several synthetic steps (Fig. 18). It is noteworthy that copper-free Huisgen cycloadditions between simple alkynes and azides were recently reported under microwave irradiation at 100  C for 30 min in the case of molecular compounds, thanks to the use of glycerol as solvent (RodríguezRodríguez et al. 2015). This method will of course be very interesting to exploit for materials functionalization. On the other hand, thiol-ene reactions are often faster than the other forms of click reactions due to the high reactivity of thiols. Thiol-ene reactions are also extremely wide in scope due to the availability of a large number of commercially available

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thiols and olefins and can proceed in the presence of organic solvents or in water, and often under mild reaction conditions. This allows its use for the selective immobilization of biological species such as proteins and enzymes without deterioration of their activity. Though the use of a metal catalyst such as Cu is eliminated in thiol-ene click reactions, it is important to note that the biomolecules used should be tolerant to the reaction conditions, especially under UV irradiation (350 nm) which is often the case for photochemically initiated thiol-ene reactions. Furthermore, the reaction conditions can also promote the formation of disulfide by-products as a result of radical recombination (Lowe 2010), while competitive reactions such as the addition of amines on acrylates can also occur. Finally, the main drawback of this reaction is the lack of obvious spectroscopic signature to easily follow the success of the reaction. Indeed, the FTIR bands of thiol, alkene, or thioester groups are weak, and they often overlap with other functional groups. Actually, the studies comparing the different click reactions on similar systems are still scarce (Göbel et al. 2016; Liu et al. 2015), and complete comparisons of the functionalization of thiol–silica, alkene–silica, or azide–silica still need to be performed for various functional groups. In a notable study, sol-gel precursors based on the phosphorylcholine moiety and synthesized either via the CuAAC or the thiol-ene click reactions (>90% yield) were click-grafted onto silica beads. It appears that both types of precursors graft easily with high loading on the surface of the silica spheres, but the loading rate appears significantly higher for the thioether linker compared to the corresponding triazole (0.63 vs. 0.44 mmol g 1), probably because of a higher steric hindrance of the triazole ring (Liu et al. 2015).

Combination of Various Click Reactions CuAAC and thiol-ene click reactions can be used in conjunction to create multifunctional hybrid materials by exploiting the advantages associated with both these click reactions as explained above. Bifunctional MSNs were synthesized by co-condensation of TEOS with two chemically orthogonal organo(triethoxy)silanes (one containing azide and the other containing a terminal olefin) (Dickschat et al. 2012, 2013). Alkyne-terminated sulfonic acid was clicked to the azide group and the olefin group was functionalized with cysteamine using thiol-ene click chemistry with AIBN as the photoinitiator. To take this one step further, the compatibility of nitroxide exchange reaction was also tested, for which bifunctional MSNs were first prepared by the co-condensation of TEOS, azide and an alkoxyamine-based precursor followed by functionalization with alkyne-terminated sulfonic acid using CuAAC. These MSNs were further functionalized to form a series of acid/base-functionalized organic/inorganic-hybrids, using the thermal nitroxide exchange reaction with different nitroxides. Moreover, bis(clickable) MSNs that can be functionalized twice by CuAAC were produced either as nanospheres or nanorods by the co-condensation method from TEOS, AzPTES, and PTESPA, with randomly distributed functionalities in controlled loadings (1%,

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Fig. 19 Alternate arrays of azide- and alkyne-bearing clickable mesoporous silica microdots and their functionalization with green and red fluorophores

2%, 5%). FRET experiments between two different clicked fluorophores evidenced that the functions are homogeneously distributed within the MSN, and that the fluorophores are separated by ca 3.5 nm for a loading of 1% of both alkyne and azide (Noureddine et al. 2015, 2016). Mesoporous silica microdot arrays produced by ink-jet printing, and featuring alternate rows of alkyne–silica and azide–silica microdots, could be produced by printing separately two different sols. These arrays were successively functionalized with two different fluorophores, with high efficiency in both cases (Fig. 19). This strategy is being applied for the production of biosensors (de los Cobos et al. 2012). Hybrid mesoporous azide–thiol bifunctionalized silica (AzSH-silica) films with vertically aligned mesopores were also produced in a one-step process using EASA, from TEOS, AzPTMS, and MPTMS in variable molar ratios. These films were then functionalized using CuAAC and thiol-ene or vice-versa. Two combinations were investigated, ethynylcobaltocenium/vinylferrocene and propargyl alcohol/allyl cyanide. The click reactions could be performed in any order without any side reactions that could potentially affect the next click reaction and the derivatization process did not affect the hexagonal mesostructure and vertical orientation of the mesopores (Vilà et al. 2016). To conclude this section, it appears that the click reactions, mainly the thiol–alkene coupling and the CuAAC reaction, can be widely used to functionalize silicas under various shapes. Apart from silicas, it is noteworthy that click reactions can also be applied to the organic functionalization of other metal-oxide materials, such as titania or iron oxide thanks to the availability of carboxylic or phosphonic acids featuring alkyne or azide terminal groups (Unger et al. 2015; Toulemon et al. 2011). This can be exemplified by the covalent anchoring of iron oxide NPs on gold substrates to form films of clustered or aligned magnetic nanoparticles (Toulemon et al. 2013, 2016).

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Technological Applications Application in Separation Science Separation processes usually feature a mobile phase running through a stationary phase, in a column which can be packed or monolithic. Due to technological constraints, the setup should feature low charge losses to allow fast kinetics and scalability. Therefore, post-functionalization methods are preferred for the development of such columns, as they will allow the possible screening of various functional groups while preserving the overall porosity and texture of the packing material. The separation efficiency relies on tunable interactions between the analytes and the stationary phase. The screening of various moieties can be realized easily if a click reaction with wide scope is performed on a parent clickable material. Recently, we showed that the interfacial properties of a series of bridged silsesquioxanes with comparable morphology could be tuned by performing CuAAC reactions on a parent clickable bridged silsesquioxane featuring alkyne functions with various organic azides. Materials featuring either strongly lipophilic, hydrophilic, or both lipophobic and hydrophobic characters could be obtained after derivatization with alkyl, PEGylated, and fluorinated chains, respectively (Noureddine et al. 2014). This example by itself illustrates the potentiality offered by click reactions for separation science.

Application in Chromatography The application of click chemistry in separation science has been the subject of a recent comprehensive review (Marechal et al. 2013). The thiol-ene and the CuAAC reactions are the most widely used, either using grafting of clickable fragments on pre-existing silica beads or monoliths, or more scarcely using co-condensation approaches. The added value of the click approaches lies in the possibility of anchoring ligands possessing various polar groups, difficult to graft by other methods (Marechal et al. 2013). The potential of the CuAAC reaction for the preparation of various HPLC packings was demonstrated in 2006. Columns featuring benzylic ethers, alcohols, or long alkyl chains could be prepared by this method (Guo et al. 2006). This led to the development of chromatography columns based on glycosilicas, where the glucose units (either monomeric or dendritic) are bonded through a triazole linker. This first study was applied for the purification of concanavalin A (Ortega-Muñoz et al. 2006). Soon after, several sugar-based columns were reported, which led to good separations of nucleosides, puric and pyrimidic bases, as well as monomeric and dimeric sugars (Guo et al. 2007). Cinchona alkaloids could also be grafted on azide–silica beads by CuAAC after the remote double bond was transformed into an alkyne function (Kacprzak et al. 2006). These chiral columns were able to separate the mandelic acid enantiomers, with an enhancement of enantiomeric recognition capacity when compared to a very similar column prepared by the thiol-ene reaction on the parent Boc-quinine

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(Kacprzak et al. 2010). Thus, the triazole linker appears as non-innocent for such separation processes. The participation of the triazole group was also highlighted in a study devoted to anion separation based on primary amine functions, where the triazole-methylamino group proved to separate better a mixture of five anions compared to the simple amino group. This very simple column can also separate a wide range of hydrophilic compounds such as nucleosides, organic acids, and bases (Liu et al. 2012). The thiol-ene reaction was also used for the grafting of cysteine residues through its -SH group on vinylsilica. This allowed to obtain a zwitterionic behavior, enabling the separation of oligosaccharides of various polymerization degree, as well as small peptides (Shen et al. 2011). Altogether, these examples highlight the high potential of both the thiol-ene and the CuAAC reactions for the preparation of chromatography columns. The mild conditions associated with the click reactions and their very wide scopes are often cited as major advantages. Though such approaches only seem interesting for highly polar groups or elaborated ligands, several studies have highlighted the non-innocent role of the triazole linkers formed during the CuAAC reaction, which may enhance the separation of the corresponding columns (Marechal et al. 2013).

Environmental Applications The click chemistry approach is being increasingly used for applications in waste treatment, where the materials are used for the selective adsorption of toxic products such as ions or antibiotics. Click approaches are indeed helpful for synthesizing materials featuring tunable organic functions, with comparable texture and morphology. This was exemplified by the group of Gao with the preparation of three clickable materials by the co-condensation method: (Gao et al. 2014a, b, 2015) SBA-15-N3, hollow silica NP HSN-N3 (15 nm voids) and PMO-N3. These materials were functionalized by various organic alkynes, which enabled to determine the best-performing groups for the adsorption of the ciprofloxacin antibiotic. Though the scaling up of azidefunctionalized materials may be problematic owing to the hazards associated with the large-scale synthesis of AzPTMS, the CuAAC reaction is indeed a very convenient conjugation strategy as it allowed the screening of different grafted molecules from a single starting platform inducing considerable speed and reproducibility. The selective removal of borate ions from aqueous solutions is another environmental challenge. These anions feature a strong affinity toward polyols, such as carbohydrates (Fried et al. 2012). Thanks to the mild conditions needed by click reactions, which are convenient for biomolecules grafting, glucopyranose was immobilized onto alkyne-functionalized MCM-41-type mesoporous silica, which allowed significant extraction capacities for borate anions in wastewater. Silica NPs coated with PEG acrylate polymers thanks to click chemistry were used for the extraction of lead(II) from aqueous solutions. Though the silica nanoparticles themselves are able to extract significant amounts of the cations, it was shown that the clicked polymer layer slightly improves the capacity of adsorption, while the kinetics are greatly enhanced (Li et al. 2012). The adsorption of

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cesium cations, which are present in nuclear wastes, was also studied. A strong binding entity for this pollutant is Prussian blue, which was originally anchored onto SBA-15 or Vycor glass pearls through the copper ions used during the CuAAC process and complexed onto the triazole ring and the hydroxyl group of the clicked propargyl alcohol (Turgis et al. 2013).

Application in Catalysis Catalysis is an important field of application for click-grafted materials, with use in all kinds of catalysis (acid–base, organometallic, enzymatic). On the one hand, the click reaction can be used to easily anchor a wide range of functional groups, such as biocatalysts, ligands, and organometallic complexes. On the other hand, it has also been developed for positioning individual catalytic entities in a controlled fashion.

Easy Grafting of Catalytic Moieties Enzymes are among the most challenging catalytic entities to be covalently grafted onto sol-gel materials. Indeed, a lot of research has been performed to entrap them non-covalently, but the covalent attachment prevents from any leaching during the purification steps and should result in more stable materials. Trypsin is an attractive enzyme as it allows cleaving proteins and ester functions. Two immobilization strategies were reported: (i) CuAAC on large-pore SBA-15-N3 and (ii) thiol-ene on a methacrylate-derived monolith. The former approach allows a high loading of 12%wt trypsin in silica, with a significant catalytic activity of about 20% relative to the free enzyme in solution for the hydrolysis of tosyl-arginine methyl ester (TAME) (Schlossbauer et al. 2008). The latter system was obtained by reduction of the disulfide bonds of trypsin followed by the radical addition to the methacrylate groups of the monolith. Provided methylene bis(acrylamide) was added as spacer during the click reaction, a similar catalytic activity was reported for the hydrolysis of TAME. The monolithic reactor was remarkably stable over time, with 88% of remaining activity after 100 cycles (Chen et al. 2011a). In the field of supported catalysis, the direct immobilization of organometallic complexes on silica substrates is also very challenging using classical methods (Monge-Marcet et al. 2011; Zamboulis et al. 2010). To this aim, azide- and isocyanate-functionalized silicas were prepared by grafting methods. The resulting functionalizable hybrid materials were modified via click chemistry (CuAAC or urethane formation) with ethynyl- and alcohol-bearing organometallic Pd pincers, with high yields and without any degradation (McDonald et al. 2009a). The palladium-containing materials were applied as Lewis acid catalysts in the double Michael addition reaction between ethyl cyanoacetate (ECA) and methyl vinyl ketone (MVK), and they importantly maintained the activity of their homogeneous analogue. It is noteworthy that the triazole linker itself was shown to slowly catalyze this Michael addition. The prepared material was recyclable over numerous cycles; the catalyst remained stable with no organometallic complex deterioration. This

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demonstrates the practical use and feasibility of this click-tethering strategy to produce catalytic material with active metal complexes. Similarly, metallated porphyrins were clicked onto azide-functionalized silica using the 2-[dimethylamino)methyl]-1-thiophenolato-copper(I) catalyst, which leads to low amounts of remnant copper in the material. The supported manganese porphyrin was evaluated for epoxidation reaction, with a preservation of a good catalytic activity, but with recyclability issues mainly arising from the obstruction of the pores by by-products (McDonald et al. 2009b). Chiral aza-bis(oxazoline) ligands were advantageously grafted onto magnetite@silica NPs (diameters of 7 nm for the core, 25 nm for the NP) (Schätz et al. 2009). Grafting and co-condensation methods were compared, showing that higher loadings are obtained using the co-condensation method. The activity of the prepared catalyst toward the desymmetrization of racemic 1,2-diols was found to be superior to the conventional analogues grafted on polymeric supports (MeOPEG, Merrifield resin) and exceeds the best results obtained until now on supported or nonsupported catalysts taking into account the yield, selectivity, and the recyclability. Proline derivatives constitute an important family of organocatalysts, which immobilization on solid supports is getting increasing attention (Ferré et al. 2016). For example, this enables the preparation of catalytic columns for continuous-flow reactions (Rodríguez-Escrich and Pericàs 2015). In particular, a proline–tetrazole was immobilized via a photoinduced thiol-ene click coupling onto thiol-grafted silica (Bortolini et al. 2012). The solid catalyst was evaluated for various asymmetric organocatalytic reactions. Indeed, similar levels of stereoselectivity between heterogeneous and homogeneous catalysts were observed for the aldol reaction of cyclohexanone with p-nitrobenzaldehyde. Furthermore, the stability of the immobilized species allowed the successful use of the catalyst in a packed-bed microreactor for environmentally benign continuous-flow synthesis of aldol products with complete conversion and long-term stability of the packing material. Furthermore, the CuAAC coupling has proven to be a straightforward method to anchor the MacMillan catalyst, which is useful for activating the Diels-Alder reaction between cyclopentadiene and cinnamaldehyde through an iminium mechanism (Ferré et al. 2016). This catalyst was immobilized onto MSNs, PMO NPs, and hollow PMO NPs (Fig. 15b) (Shi et al. 2011; Porta et al. 2014). Interestingly, the grafting on MSNs through a triazole linker gave similar activity but significantly higher enantioselectivity than the corresponding catalyst grafted through an ether linker on silica beads (Porta et al. 2014). Furthermore, in the case of the PMO NPs, the hollow structure favors faster conversions while the stereoselectivities are higher in the case of the click-grafted material (initially prepared by co-condensation) compared with the one obtained by conventional post-functionalization (Shi et al. 2011).

Precise Positioning of Catalytic Functions Some organometallic catalyzed reactions require the formation of 1:1 ligand: metal adducts, or may be slowed down by the formation of dimetallic complexes. Though

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these limitations may be overcome in solution by dilution effects, the immobilization of such catalytic systems is more challenging, as site-isolated functions are required (Conley et al. 2014). The formation of clickable materials with low loading of randomly distributed derivatization sites is therefore a good strategy to manage materials with longer average inter site distances while managing high enough catalyst loading (Fig. 10). This was exemplified on a series of SBA-15-N3 materials synthesized by a co-condensation method between TEOS and AzPTES with controlled loadings of organosilane (x = 0.2–8 M %TEOS) (Nakazawa et al. 2012). Though the catalytic effect of a clicked TPA-Mn2+ (TPA: tris(pyridylmethyl)amine) system in the epoxidation of oct-1-ene was found to be closely related to the surface coverage, with maximum activity for site-isolated ligands, the activity of FeTPP complex for a carbene transfer reaction remains independent of the amount of grafting sites, as the porphyrin complex is particularly stable. A related Ni-TPA complex was investigated for the oxidation of cyclohexane into cyclohexanol (Nakazawa et al. 2013). For such complexes, the oxidation reaction is favored for a 1:1 TPA:Ni complex, as opposed to the 2:1 adduct which is rather inactive. Therefore, the complexation of nickel(II) acetate onto SBA-15-N3 click-grafted with TPA ligands results in much higher catalytic efficiency when the initial loading is low. Cooperativity is also a very important phenomenon in catalyzed reactions. Indeed two different functions may act synergistically and concomitantly during the catalytic cycle. For example, the presence of both acidic and basic sites can enhance the electrophilicity of a reactant while activating a nucleophile during aldol or Henry reactions. This fact has been widely studied for supported bases on silica where the surface silanol groups synergistically aid the basic catalysts (Sharma and Asefa 2007). Research is now being performed to precisely locate two distinct functions able to concomitantly participate in a catalytic reaction. Heterogeneous cooperative catalysis on a solid support requires the control of the concentration and the arrangement/distance of the functions to activate each of them individually in a cooperative way without any destructive interference. To this aim, Eckert and Studer prepared bifunctionalizable MCM-41-type MSNs as a cooperative catalysis platform (Dickschat et al. 2012). Due to the versatility and the exceptional group tolerance of the click chemistry, azide/alkene or azide/alkoxyamine-bearing MSNs were readily addressed by orthogonal click grafting of specific catalytic functions. Sulfonic acid and amine groups with, respectively, alkyne and thiol moieties were clicked on the azide/alkene MSN in order to give an acid/base catalytic material via CuAAC and thiol-ene coupling. The comparison of the catalytic activity of the bifunctional system with a combination of the two corresponding monofunctional MSNs (grafted solely by sulfonic acid or amine) showed an important difference: while no reaction progress or important side products occur in the case of the monofunctional MSNs as catalysts, excellent yields and selectivities were obtained when both functions are grafted (Fig. 20). Taking into account the robustness of the click grafting strategy, other species were envisioned to be tethered for the cooperative catalysis of Tsuji-Trost allylation where a Pd-complex and a basic site are both needed to catalyze the reaction

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Fig. 20 Bifunctional MCM-41 for cooperative catalysis prepared from bis(clickable) nanoplatforms (Adapted from Dickschat et al. 2012 with permission from Wiley)

(Dickschat et al. 2013). Based on a parent MSN-bearing azide and alkoxyamine, it was possible to graft the 2-ethynylpyridine serving as a Pd ligand as well as different amine moieties for basic activation. Here again, the cooperativity of the catalysis between the click-grafted moieties was demonstrated.

Application in Biology Silica NPs are increasingly used for diagnosis and therapy of diseases, in particular cancer. The controlled functionalization of silica NPs is a key step for the preparation of active, stimuli-responsive nanosystems. Functionalized NPs are being studied for applications such as antimicrobial coatings, sensing of biomolecules, imaging, drug delivery, and photodynamic therapy.

Antimicrobial Surfaces Antimicrobial surfaces find many applications for clinical use, i.e., for catheters, surgical devices, or textiles. Important features are the adhesion of the biocidal coating on the substrates, and the resistance to sterilization processes. To this aim, antifouling coatings based on silver nanoparticles were recently developed through the covalent attachment by the thiol-ene click reaction of PMO NPs featuring sulfonic acid and thiol groups on glass substrates, followed by silver NP growth (Gehring et al. 2015). This resulted in glass coatings with strong biocidal properties toward Pseudomonas aeruginosa, with cooperative effects between the sulfonic acid groups and the AgNPs, the superacid function helping to release Ag+ cations from the NPs. The cationic polymer poly(L-lysine) also exhibits strong antimicrobial activity. This polymer was attached onto silica NPs (20 nm) using the thiol-ene coupling,

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resulting in silica@polymer NPs, which were used as antimicrobial agent toward gram-positive Bacillus subtilis and gram-negative E. coli (Kar et al. 2010).

Sensing The click reaction is particularly suited to anchor biomolecules or analogues onto the surface of silica materials for sensing applications. The CuAAC reaction has been used to anchor a Fe–biuret complex that mimics horseradish peroxidase and used for glucose detection in the presence of glucose oxidase in physiological conditions. Interestingly, this system was tested in real biological samples, such as mice blood plasma (Malvi et al. 2012). A different approach has been developed for the concentration and titration of thrombin, a protein associated with several diseases such as pulmonary metastasis or synovial inflammation (Wang et al. 2015). An aptamer-derived monolithic column was prepared using the thiol-ene click reaction between a vinyl-functionalized monolith and a thiol-derived DNA strand. The thrombin accumulates in the column, with ca 92% extraction recovery, and can then be assayed by the spectrophotometric titration of the hydrolysis of a peptide p-nitroanilide. In vitro imaging of sialic acid, which is overexpressed at the surface of HeLa cervical cancer cells, was realized using phenylboronic acid-tagged silica NPs doped with fluorescein (Cheng et al. 2013). The phenylboronic was linked through a thiolene coupling reaction. Cancer Therapy MSNs are being intensively used for drug delivery in cancer applications (Giret et al. 2015). MSNs of ca 100–130 nm are preferred, as they may accumulate efficiently in cancer tissues through the EPR effect (Lu et al. 2010). Hydrophobic drugs, which are usually difficult to administer, can be easily encapsulated in MSNs and delivered to the targeted tissues avoiding the use of high doses of expensive and toxic drugs. The MSNs can be passively endocytosed, though their functionalization with targeting moieties enhances the penetration rate (Porta et al. 2013). The release of the active principle should occur preferentially inside the cells, while the cargo should be kept within the nanovehicle outside the target. To achieve this goal, poregating mechanisms such as nanovalves have been developed and anchored at the external surface of MSNs. The construction of such gatekeepers while preserving the cargo inside of the pores requires performing reactions under mild conditions (low temperature, in water). Click reactions such as the CuAAC and the thiol-ene couplings are thus widely employed to this purpose. The construction of such nanomachines employing the CuAAC reaction was first demonstrated by the groups of Zink and Stoddart (Patel et al. 2008). Snap-Tops featuring bulky adamantyl groups and linked through ester or amide functions were built on MSNs by the CuAAC reaction in the presence of rhodamine B in the pores and adamantyl groups around triethylene glycol stalks at low temperature and in water (Fig. 21). The nanocontainers featuring the ester function could be opened in the presence of porcine liver esterase to release their cargo, while those containing the amide groups did not exhibit cargo release. The same strategy was later

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Fig. 21 Different steps for the construction of the first nanomachine using CuAAC as a derivatization strategy (Adapted from Patel et al. 2008 with permission from the American Chemical Society)

employed on linkers built on disulfide groups upon activation with thiols (Ambrogio et al. 2010). Peptide-based enzyme-responsive nanosystems were developed soon after by grafting alkyne-containing peptides onto N3-MSNs. These systems were cleavable in the presence of proteases, while no release of cargo was observed at neutral and slightly acidic pH in the absence of such enzymes (Coll et al. 2011). Other types of gate-keeping mechanisms that can be anchored by the CuAAC reaction include selfcomplementary DNA duplexes. In this case, after the DNA strands were clickgrafted, a denaturation-reassembly process was needed to achieve an efficient capping (Chen et al. 2011c). The pores could be opened either upon heating or in the presence of endonucleases. This enabled the release of camptothecin in vitro on HepG2 human liver cancer cells. More sophisticated nanomachines make use of an antenna (donor) transferring energy, electrons, or protons to an actuator (acceptor) (Croissant et al. 2013, 2014a; Guardado-Alvarez et al. 2014). In this case, the precise relative position of the donor and the acceptor may be important to control in order to obtain optimized properties. This is particularly true for nanomachines based on proton transfer. In a first study, the photoacid (donor) and the pH-triggered nanovalve (acceptor) were anchored by conventional grafting (Guardado-Alvarez et al. 2014). This nanomachine was efficient in water, but no cargo release was observed in buffered solutions. In a second study, a similar system was constructed by click chemistry from parent bis(clickable) nanoparticles obtained by the co-condensation of TEOS, AzPTES, and the alkynecontaining organosilane 2 (Fig. 8a) (Noureddine et al. 2015). After two consecutive click reactions to anchor a photoacid and a thread, the NPs were loaded with a cargo, then the pH-sensitive supramolecular nanogate was formed with β-cyclodextrin

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Fig. 22 Successful use of bisclicked MSN as a light-triggered proton-transfer nanomachine in buffered solutions (Adapted from Noureddine et al. 2015 with permission from the Royal Society of Chemistry)

(Fig. 22). This bifunctional nanomachine was able to release its cargo upon irradiation, even in concentrated buffer solution, which enabled its application in vivo. Short distances between the donor (photoacid) and the acceptor (basic valve) are particularly required in this case, as the targeted proton transfers should occur much faster than the acid–base reactions with the phosphate ions present in the PBS medium. Indeed, the random positioning of the clickable fragments by the co-condensation approach at low loadings is in strong contrast to the well-described aggregation of functional groups occurring with the classical grafting reactions (Hoffmann et al. 2006; Kickelbick 2004; Nakazawa et al. 2012). Apart from drug delivery, cancer treatment can also involve photodynamic therapy. This process is based on the production of singlet oxygen or reactive oxygen species upon irradiation of a photosensitizer. In a recent example, a high molecular weight photosensitizer was derivatized under anhydrous conditions with four triethoxysilyl groups through triazole moieties by CuAAC (Croissant et al. 2015b). This compound could be transformed into discrete bridged silsesquioxane NPs that may be activated upon two-photon excitation in the near infrared. The obtained nanoparticles were strongly toxic under irradiation toward MCF-7 breast cancer cells. These phototherapic properties could be advantageously combined with drug delivery for PMO nanoparticles based on ethenylene bridging fragments and containing the photosensitizer. In this case, a synergistic effect of the phototherapy and the doxorubicin delivery was observed (Croissant et al. 2014b). Overall, this section demonstrates that the rapid expansion of nanomedicine by MSNs over the past decade has benefited from the simultaneous development of the click methodologies, which not only simplify the anchoring of fragile species under mild conditions but also provide new access to elaborated organosilane precursors and enable more precise positioning of functional moieties.

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Conclusion Among the post-functionalization methods, the thiol-ene and the CuAAC click reactions have been shown to be very valuable for a wide range of applications. With respect to the conventional grafting of organo(alkoxy)silanes or organo(chloro) silanes on preformed silicas, they clearly feature a much wider scope, enabling facile derivatization with strongly polar molecules such as DNA strands, peptides or sugars, as well as any kind of hydrophobic functionalities. The mild conditions associated with those reactions have allowed to easily graft such sensitive biomolecules, but also to perform the grafting while preserving sensitive fragments at the surface of the material, or cargos within the pores. Thanks to the high conversions observed with such click reactions, any clickable fragment may be transformed into a valuable functional group. On the one hand, when high loadings of clickable fragments are present in the starting material, the functionalization leads to massive changes in the surface properties, enabling facile screening of surface modifiers while preserving the morphology for applications like chromatography or extraction. On the other hand, at low loadings of clickable functions, and when the parent material is synthesized by co-condensation of a silica precursor with a non-selfassembling clickable organosilane, the clicked functions are mostly site-isolated and randomly distributed within the whole material. In particular, this is very interesting for supporting catalysts on silica supports. By extending this concept to materials obtained by co-condensation and featuring two or more fragments that can be transformed through orthogonal click reactions, it is possible to construct elaborated nanosystems with communicating functions, randomly distributed over the surface. This was indeed applied to cooperative catalysis and nanomedicine. Finally, the increasing application of click chemistry to silica materials benefits from the simultaneous advances in sol-gel technologies, with the preparation of mesoporous nanoparticles, microdot arrays, or thin films featuring oriented pores that have been developed over the past 15 years. By easily linking functional (bio) organic fragments to original inorganic scaffolds, click methodologies applied to solgel materials pave the way to the preparation of advanced materials with improved properties for separation processes, sensing, catalysis, or medicine.

References Achatz DE, Heiligtag FJ, Li X, Link M, Wolfbeis OS. Colloidal silica nanoparticles for use in click chemistry-based conjugations and fluorescent affinity assays. Sens Actuators B. 2010;150 (1):211–9. https://doi.org/10.1016/j.snb.2010.07.014. Ambrogio MW, Pecorelli TA, Patel K, Khashab NM, Trabolsi A, Khatib HA, Botros YY, Zink JI, Stoddart JF. Snap-top nanocarriers. Org Lett. 2010;12(15):3304–7. https://doi.org/10.1021/ ol101286a. Baskin JM, Prescher JA, Laughlin ST, Agard NJ, Chang PV, Miller IA, Lo A, Codelli JA, Bertozzi CR. Copper-free click chemistry for dynamic in vivo imaging. Proc Natl Acad Sci U S A. 2007;104(43):16793–7. https://doi.org/10.1073/pnas.0707090104.

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Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Synthesis of Sol-Gel Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . “In Situ” and “Ex Situ” Methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nanocomposites with Oxide Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nanocomposites with Metal Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nanocomposites with Chalcogenide or Halide Nanoparticles . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nanocomposites with Polymer Matrix . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Applications of Sol-Gel Nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Biomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Chemical Sensors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Catalysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Wettability . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Scratch-Resistant Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

This chapter describes nanocomposites where at least one of the phases has been obtained by the sol-gel method. In the first part, the synthesis of the nanocomposites is considered, and an attempt is made to resume the general principles and methods related to different classes of materials, with oxide, hybrid organic–inorganic, and polymeric matrices and oxide, metal, and non-oxide nanoparticles. The second part is devoted to the applications of sol-gel nanocomposites. Representative examples of nanocomposites that potentially can have commercial application are described, based on the literature of the last 10 years.

M. Guglielmi (*) · A. Martucci Dipartimento di Ingegneria Industriale, Università di Padova, Padova, Italy e-mail: [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_100

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Introduction The development of sol-gel nanocomposite materials has been mainly driven by their potential use in many applications, and commercial sol-gel nanocomposites are already part of manufacturing technology. IUPAC defines a nanocomposite as a “composite in which at least one of the phases has at least one dimension of the order of nanometers,” where “composite” is a “multicomponent material comprising different (nongaseous) phase domains in which at least one type of phase domain is a continuous phase.” We may define sol-gel nanocomposites as those in which at least one phase is obtained by the sol-gel method. Often, in the literature, homogeneous monophasic hybrid sol-gel materials are named nanocomposites, in clear contradiction with the above definition. Of course, borderline situations exist, where the distinction might be a matter of discussion, but these will not be considered here. The matrix (continuous phase) material, usually classifies composites, allowing to distinguish among metal, ceramic, polymer matrix. In the case of sol-gel nanocomposites, metal matrix is excluded, but a further type of matrix, the hybrid organic–inorganic one, is added. As for usual composites, the combination of different materials allows to obtain a new material with new and tailored properties (mechanical, thermal, electrical, or optical). The most important difference between a traditional composite and a nanocomposite is the area of the interface between the matrix and the dispersed phase. In nanocomposites, this is much larger, and very often this allows to obtain huge effects even with small amount of the second phase. Furthermore, the specific properties of the nanophase are introduced. The application of the sol-gel method to the synthesis of nanocomposite materials started to emerge in the literature approximately at the end of the 1980s. Since then the number of published papers increased continuously. Looking at the number of published papers, as reported by the Web of Knowledge using the keywords “sol-gel” and “nanocomposites,” it comes out that from the beginning of the millennium the growth has been linear, arriving at more than 500 in 2015. Initially the idea to start from a single batch was very attractive, and the first efforts were addressed to the so-called “in situ” strategies. The development of hybrid organic–inorganic materials was almost contemporary and gave an important contribution to the control of the matrix-dispersed phase interface. Then “ex situ” methods started to attract the interest of researchers because of the possibility to better control the final microstructure of the nanocomposite. Among the most interesting features of the sol-gel approach to nanocomposites is the possibility to get them as thin films to coat every type of substrates, and many applications deal indeed with high-tech coatings and varnishes.

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Synthesis of Sol-Gel Nanocomposites The synthesis of nanocomposites is successful if one succeeds to get a homogeneous dispersion of nanometer size particles, constituting the dispersed phase, in the matrix, avoiding aggregation and particle growth phenomena. This is not simple at all. Aggregation phenomena are driven by the tendency of the system to reduce the interface energy, which is very high in the case of nanoparticles. When the nanocomposite is obtained from a single batch where nanoparticles are dispersed in the sol which will originate the matrix, the dispersions may be stabilized electrochemically (thermodynamic approach) by increasing the electric double layer thickness or through steric stabilization (kinetic approach) by decorating the surface with surfactants. However, aggregation may occur also in the already-synthesized nanocomposite. The stability of the nanostructure with time and temperature depends on the tendency, and possibility, of the nanoparticles to aggregate. There are two different strategies to try to solve this problem: modifying the surface of the nanoparticles to reduce the interface energy or reducing the mobility of the particles, for example, by linking them to the matrix with covalent bonds. This last approach is largely used in the synthesis of sol-gel nanocomposites. The interface between the matrix and the dispersed phase is important not only for avoiding or limiting aggregation and growth but also for controlling the properties of the nanocomposite material. Nanoparticles can be functionalized using monofunctional or bifunctional ligands. In the first case, the effect is limited to the modification of the nanoparticle surface, leaving the control of interface reactions to weak interactions with the matrix. Bifunctional ligands are, instead, the way to get a strong interaction with the matrix. The nature of the ligand determines the type of interaction and the effect on specific properties. For example, linking the nanoparticles to the matrix network helps to avoid (or limit) nanoparticle aggregation; using spacer groups between the two reacting groups in the ligand and changing their length and flexibility allow to control the stress transfer between the matrix and the particle, tailoring the mechanical behavior (stress relaxation capability, fracture toughness, elastic moduli, hardness, resilience modulus, etc.). There are three main general approaches used to modify the surface of nanoparticles: chemical treatments, grafting of polymers, ligand exchange (for a review, see Kango et al. 2013). Chemical treatments are those where the surface of nanoparticles are modified using moieties, different from the native ones. As an example, both oxides and metals, which surface is usually oxidized, exhibit hydroxyl groups which may react with alkoxysilanes, epoxides, and alkyl or aryl isocyanates. Inorganic nanoparticles may be “decorated” with polymers by the “grafting” approach. There are two different strategies, “grafting to” and “grafting from,” schematically represented in Fig. 1 (Achilleos and Vamvakaki 2010). The “grafting to” method is based on linking macromolecules onto the surface of the nanoparticle using the affinity of the polymer end group for the functionalities present on the

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Fig. 1 Schematic representation of the “grafting to” and “grafting from” methods for the decoration of particles with polymers

Fig. 2 Schematic representation of the ligand exchange method

particles. The “grafting from” method is, instead, based on the growth of the polymer chains from initiators present or generated on the particle surface. This second approach allows to obtain, in general, a higher density of grafted molecules. Ligand exchange is a very important reaction for preparing functionalized inorganic nanoparticles. It is based on the substitution of ligands present on the particle surface (outgoing ligands) with ligands present in the surrounding environment (incoming ligands), which link to the surface stronger than the substituted ones (Fig. 2). This method makes it possible to modify nanoparticles with ligands that cannot be introduced during their synthesis. Although the comprehension of the mechanisms is not complete, the chemistry offers many opportunities, and this is a very important method to design both composition and structure of nanoparticles surfaces (Caragheorgheopol and Chechik 2008). Playing with sol-gel allows to work also on the matrix to optimize the interaction with nanoparticles, by introducing functional groups in the network. The development of hybrid organic–inorganic materials opened many new possibilities in this direction.

“In Situ” and “Ex Situ” Methods Sol-gel nanocomposites may be obtained following “in situ” or “ex situ” strategies. In situ methods are based on the preparation of one batch containing the precursors for the matrix and the dispersed phase. Nanoparticles are formed by chemical reaction contemporary to the network development or by thermally induced phase separation in

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the already generated matrix (inorganic or hybrid organic–inorganic) in the case of “all sol-gel” nanocomposites or in a polymer matrix through sol-gel reactions. An advantage of the in situ approach is the reduced number of steps in the preparation. A major drawback is given by the possibility of interactions between the processes giving rise to the two (or more) phases, which makes the control in the development of the micro-/nanostructure very difficult or sometimes even impossible. The ex situ methods are based, in general, on the separation of the procedures to get the matrix and the dispersed phases, and the systems are mixed together in a second step. The separation allows to get a much finer control of the final microstructure by avoiding as much as possible any interactions among different chemical processes. The “in situ” approaches for the synthesis of sol-gel matrix nanocomposites take advantage of the potentiality of the sol-gel method to include, in one liquid batch, all the precursors necessary for the development of the matrix and the dispersed phase. In general, the matrix is obtained by the usual hydrolysis and condensation reactions, and the precursors for the dispersed phase remains inside the matrix as precipitated salts or in a dispersed ionic state. In the last case, the phase separation is induced by chemical reactions or by thermal treatment. A general scheme is reported in Fig. 3. A sol containing the precursors for the matrix and the dispersed phase is prepared starting, generally, from two separate solutions. Both may be, in some way, processed before mixing, but the result of mixing is a one-phase system. Hydrolysis and condensation of the precursors, usually alkoxides, give rise to a gel matrix that may be transformed in a xerogel or an aerogel upon drying. The dispersed phase may be generated in the liquid sol, in the wet gel, or in the dried xerogel or aerogel by means of chemical reaction between metal ions and anions provided by components already present in the original sol or introduced in a later stage. A second way to get the dispersed phase is by thermal treatment of the gel, during the drying process or during densification at higher temperatures. DP precursor(s)

Matrix precursor(s)

Sol Gel Thermal treatment for DP formation

Chemical reaction for DP formation

Xerogel/Aerogel

Final nanocomposite material

Fig. 3 General scheme of the “in situ” synthesis strategy. DP dispersed phase

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The in situ synthesis has been used mainly to prepare nanocomposites made of an oxide or hybrid organic–inorganic matrix, but there are interesting examples of polymeric matrix nanocomposites with oxide nanoparticles generated “in situ” by the sol-gel method. The aim of the “ex situ” approach is to keep the processes for the production of nanoparticles and for the synthesis of the matrix independent as much as possible, in order to reduce the interactions between the two processes. Clearly ex situ processes may be the combination of different synthesis methods. It will be referred to sol-gel nanocomposite if at least one of the phases is obtained through sol-gel reactions. In the case of an “all sol-gel” nanocomposite, the general scheme for the synthesis is the one shown in Fig. 4.

Nanocomposites with Oxide Nanoparticles In the case of oxide nanoparticles dispersed in a sol-gel inorganic matrix, the chemical affinity is much higher and there is competition between self-aggregation of particles and linking (through weak or strong bonds) of particles to the matrix, which might, in some cases, avoid the agglomeration problem. With hybrid organic–inorganic matrices, the nature of the organic moieties may modulate the affinity, allowing a better control of the interaction. The results reported in the literature seem to confirm this picture, offering examples of homogeneous nanocomposites obtained in very simple ways, without specific chemical modifications, and examples where a quite complicated chemistry is necessary to get acceptable results. Fig. 4 General scheme of the “all sol–gel” “ex situ” synthesis strategy. DP dispersed phase

Matrix precursor(s)

DP precursor(s)

Matrix sol

DP sol

Gel

Xerogel/Aerogel

Final nanocomposite material

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Looking at “in situ” methods, important basic concepts have been published by Breitscheidel et al. (1991) in a paper devoted to the preparation of metal complexes in inorganic matrices. As a first step of the synthesis, the authors obtained a dispersed metal oxide phase by using bifunctional ligands of the type X – Si(OR)3, with X = NH2, NHCH2CH2NH2, CN, CH(COMe)2, able to form stable complexes with transition metals and, at the same time, to participate to the sol-gel reactions. This approach was the key to get a homogeneous distribution of the metal ions into the gel network. After thermal removal of the organic groups, oxidation of the metal ions took place, giving oxide nanoparticles. Several examples of the application of this concept may be found in the literature. Our group published several papers, since 2003 (Martucci et al. 2003a), with the goal to fabricate optical gas sensors sensitive to reducing gases (CO, H2, etc.). The nanocomposites were based on sensitive oxide nanoparticles, such as NiO, SnO2, and Co3O4, finely dispersed in a porous silica matrix. Up to 40% mol metal oxide could be homogeneously distributed in the matrix by using NH2(CH2)3Si (OR)3 (aminopropyltriethoxysilane, APTES) or N-[3-(trimethoxysilyl)propyl]ethylendiamine (DAEPTMS) as bifunctional ligands and a mixture of TEOS and MTES for the matrix (Martucci et al. 2003b, 2004a, b; Cantalini et al. 2005) (Fig. 5). The same route was used to obtain more complex nanocomposite films, containing both NiO and Au nanoparticles in silica (Buso et al. 2007a) and in titania (Della Gaspera et al. 2010a) matrices. With “ex situ” methods, the most important aspect for the successful synthesis of homogeneously dispersed nanocomposites is Fig. 5 TEM image of SiO2 porous matrix containing 20% mol of NiO nanoparticles (From J Am Cer Soc, Nanostructured Silicon Oxide–Nickel Oxide Sol‐Gel Films with Enhanced Optical Carbon Monoxide Gas Sensitivity, 86, 2003b, 1638–1640, Martucci A, Pasquale M, Guglielmi M, Post M, Pivin JC, Springer Science+Business Media with permission of Springer)

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the surface modification of oxide nanoparticles. Due to hydrogen bonds, the particleto-particle interaction is very strong, and it is very easy to get irreversible agglomeration when the distance among particles decreases with change of pH or evaporation of the solvent, as clearly discussed by Schmidt et al. (1998). An effective way to keep the interaction under control is the substitution of hydroxyl groups with controllable functional moieties with tailored reactivity as previously discussed. An example can be found in Peeters (2000) where colloidal silica particles have been incorporated in silica or silica-based hybrid organic–inorganic sol-gel matrices by functionalizing the particles with methyltrimethoxysilane (MeTMS) and using the same molecule as precursor for the matrix. Critical for the success of the synthesis was the size of the oligomers formed by condensation of the hydrolyzed species generated during the pre-hydrolysis step. Large oligomers were found to not react with the silica surface, probably due to their apolarity, while small oligomers could cover the silica particles with approximately a monolayer of MeTMS molecules, with their methyl groups probably oriented toward the outside. A second interesting example is given by the use of boehmite as nanofiller for hybrid matrices. Commercially available boehmite nanoparticles could be stabilized in a mixture of GPTS and TEOS, by substitution of acetic acid, the initial stabilizer in water, with the silanes (Schmidt et al. 1998). In this case, a further interaction between particles and the matrix was given by the influence of Al-OH groups, behaving as Lewis acid, on the epoxy ring opening and organic polymerization.

Nanocomposites with Metal Nanoparticles Nanocomposites with metal nanoparticles are of interest because of their optical, electrical, and catalytic properties. The in situ synthesis of this type of nanocomposites follows the general scheme reported in Fig. 6. All the precursors necessary for the formation of the matrix and the dispersed phase are mixed, and the formation of metal nanoparticles may occur by different reduction processes directly in the sol or in the gel, at different stages of its drying and thermal treatment. Nanocomposites containing a metal oxide as dispersed phase may be transformed by reduction of MO nanoparticles into metal nanoparticles. Sodium citrate, hydrogen peroxide, citric acid, hydrogen, formaldehyde, etc., may be used as reduction agents, along with polymer stabilizers, such as polyvinyl alcohol and sodium polyacrylate, to control the particle growth. Both the reduction agent and the stabilizer have an important effect on the size and size distribution of particles. Stronger reduction agents favor a fast formation of many nuclei, giving smaller particles. With slow reduction reactions, both wide and narrow size distributions can be obtained, depending on parameters such as the concentration of precursors, pH, temperature (Cao 2004). Three different approaches have been reported in the literature in early 1990s, putting the basis for further development: (a) reduction of metal oxide to metal

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Fig. 6 General scheme for the “in situ” synthesis of nanocomposites based on sol–gel matrices and metal nanoparticles

nanoparticles by thermal treatment under hydrogen (Breitscheidel et al. 1991), (b) treatment of an aged gel with a reducing solution (i.e., formaldehyde to completely reduce silver) (Datta and Das 1992), and (c) reduction of a metal salt in the sol by UV irradiation at room temperature (Mennig et al. 1992). The important basic ideas in these reports are as follows: the reduction of metal ions may be achieved at different steps of the process; the reduction may be induced by chemical reactions, irradiation, or thermal treatment; and the control of growth and dispersion by linking of metal ions or complexes to the originating matrix network may be obtained by the use of multifunctional silanes, which form complexes with unreduced metal ions at the surface of clusters. The thermal precipitation in the xerogel at relatively high temperatures is certainly the worst from the point of view of size and size distribution control. Furthermore, metal ions may strongly affect the pore formation process of silica gel during drying and heating (Takahashi et al. 2000). The generation of particles in the sol is, instead, the most promising approach, especially if the sol-gel process for the matrix formation may be, in some way, taken

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separated from the one for NPs growth. This idea is at the base of the “ex situ” methods, but brought to an interesting intermediate approach proposed by Giacoin (Giacoin et al. 1994). Silver and gold NPs were precipitated at room temperature in the sol containing the metal precursors (AgNO3, HAuCl4) by irradiation with γ-rays. Hydrolysis and condensation reactions of alkoxides were then allowed by slow diffusion of water from a H2O-saturated outside inert atmosphere. The growth of metal nanoparticles is obtained by colloidal chemistry in the same solution that will give rise to the matrix, but the two steps are sequential in time, allowing a better control of both. An “ex situ” method to prepare Au-loaded TiO2 thin films was reported by Della Gaspera et al. (2012). It is based on mixing 5% of gold nanoparticles, synthesized by the Turkevich method, and 95% of titania nanoparticles, obtained by hydrolysis of titanium isopropoxide, followed by thermal sintering. Highly crystalline porous nanocomposite films sensitive to CO were obtained by this way.

Nanocomposites with Chalcogenide or Halide Nanoparticles Silicate glasses doped with chalcogenides or halides nanocrystals were firstly obtained by the sol-gel method by Nogami in the early 1990s (Nogami et al. 1990a, b, 1991a, b, c, 1992, 1994). An “in situ” synthesis was used to get Cd- or Pb-doped xerogel, and metal sulfide nanocrystals were obtained by reaction with H2S at room temperature. The sulfide nucleation and growth were possible because of the high porosity of the gel, which allowed the reacting gas to react with metal ions dispersed in the silica network. A different approach was based on the addition of the anion precursors in the starting sols, and by precipitating the dispersed phase upon thermal treatment in a reducing atmosphere, to avoid oxidation of the chalcogenides (Tohge et al. 1992). A disadvantage of both these methods is that the porous xerogel exposes the non-oxide NPs to oxygen and, therefore, to oxidation. On the other hand, the densification achievable at temperatures higher than 1,000  C leads to transformation of the nanoparticles, both in size and composition. A solution to this problem is the use of matrices able to sinter at lower temperatures, as demonstrated by Takada et al. (1992) who used a borosilicate composition able to fully densify below 600  C. Also hybrid organic–inorganic sol-gel matrices with lower porosity may have a similar effect (Guglielmi et al. 1994). An interesting nanocomposite material based on a hybrid organic–inorganic ureasilicate matrix and CdS nanoparticles was prepared by Boev et al. (2004) by the formation of amorphous clusters of CdS, through reaction of cadmium nitrate with thioacetamide, C2H5NS, in a ureasilicate sol. The hybrid matrix was obtained using a organically modified silicon alkoxide and a diamine functionalized oligopolyoxyethylene, which formed a rubberlike material composed of a siliceous network grafted through urethane groups to both ends of poly(oxyethylene) glycol segments (Goncalves et al. 2002). The final material, which didn’t require any heat treatment, was transparent, homogeneous and flexible, and dense enough to prevent further aggregation and sedimentation of the CdS colloidal phase.

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A different strategy to protect the non-oxide nanocrystals from oxidation is to fill the pores by further impregnation of the porous xerogel with the sol (Piñero et al. 1994; Litrán et al. 1997), but it was found that the impregnation process influences the xerogel evolution, becoming an important step for the control of the final microstructure (Hummel et al. 1997). The examples given above show that the “in situ” approach is useful and relatively simple, but the interference of various chemical processes occurring at the same time in the same batch is difficult to control, especially with respect to the size and size distribution of nanoparticles. An “ex situ” approach, based on the addition of already-synthesized NPs to the sol, was proposed to avoid these problems (Spanhel et al. 1992). An example is given by the synthesis of CdS or PbS quantum dots by reaction of Cd or Pb acetates with thioacetamide in presence of a mercaptosilane and the subsequent addition of the colloidal sol to the matrix precursor sol (Guglielmi et al. 1997; Martucci et al. 1999a, b). A homogeneous distribution in the matrix, with controllable particle diameter between 2 and 5 nm, and narrow size distribution were easily obtained (Fig. 7). This approach is very flexible and allows the synthesis of complex multifunction systems. An example has been experienced by the authors with the preparation of a material which was required to have high transparency in the visible range, tunable refractive index, luminescent properties, and being suitable for imprinting lithography and aerosol printing technologies (Antonello et al. 2011). The problem was afforded by designing a hybrid organic–inorganic matrix containing titanate nanosheets to increase the refractive index and CdSe@ZnS core-shell quantum dots to tune the photoluminescence properties. The two types of nanoparticles

Frequency (%)

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Fig. 7 HRTEM micrograph of a zirconia ormosil film doped with 10 mol% PbS and histogram of PbS particle size distribution (Reprinted from J Non-Cryst Solids, 244, Martucci A, Innocenzi P, Fick J, Mackenzie JD, Zirconia-ormosil films doped with PbS quantum dots, 1999b, with permission from Elsevier)

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were synthesized separately following specific protocols, but in both cases 6-amino1-hexanol was used to functionalize the surface of the nanoparticles and to allow them to link to the matrix network.

Nanocomposites with Polymer Matrix Nanocomposites with a polymer matrix, where sol-gel has been used to obtain oxide nanoparticles dispersed in the macromolecular organic network, are of great interest because their mechanical properties may be significantly improved. It is well known that in traditional composites, the mechanical properties (i.e., the elastic modulus) are modified by the filler. With nanofiller it should be possible, in principle, to get much higher effects with smaller fraction due to the exceptionally high surface to volume ratio of nanoparticles and/or, in some cases, the extremely high aspect ratio. This would be true, however, only in the case of a homogeneous dispersion of the dispersed phase in the polymer and a good interaction between the two phases. Therefore, the molecular control of the metal oxide/polymer interface is the key for success. A possible “in situ” strategy is based on the contemporary occurrence of hydrolysis and condensation reactions of the inorganic precursors and the organic polymerization. The possible interactions between the two processes have been described by Kickelbick (2008). They may be summarized as follows: (a) The alcohol molecules present in the system or produced in the hydrolysis of alkoxides may behave as chain terminators in anionic polymerizations. (b) Initiators and chain ends may act as nucleophile species and lead to substitution reactions at the metal alkoxide, resulting in chain termination. (c) Some groups present in polymers (i.e., ethylene oxide in PEO) interacts with the water molecules necessary for the hydrolysis in the sol-gel process, thus affecting the kinetics of gelation. The compatibility between the inorganic phase and the polymeric one may be controlled in different ways, for example, using “compatibilizing” solvents. An example is given by N,N-dimethylformamide or N,N-dimethylacetamide that forms hydrogen bonds with silanol groups of the silica network and is highly compatible to organic polymers. Polymers can also be modified by functionalizing their ends or the backbone with groups that increase the interaction between the organic and the inorganic components. Depending on the effectiveness of the interaction, it is possible to drive the process to give a homogeneous hybrid system or a nanocomposite. Due to the interactions among the different chemical species in the system, the process must be optimized, also from the point of view of the sequence of addition of the reactants. Hirano and co-workers demonstrated that if the organic polymerization precedes the hydrolysis and condensation reactions, the composite will be made of small

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particles, but if the hydrolysis is run before the polymerization, larger particles are obtained (Yogo et al. 1994; Hirano et al. 2001, 2003). Many papers have been published on the synthesis of polyimide (PI)/SiO2 nanocomposites. This is a special case, indeed, as different aspects make the synthesis relatively simple. Polyamic acid (PAA), which is the precursor used for the polymer matrix, can tolerate the water necessary for the sol-gel reactions, and the imidization reaction, which converts PAA to a polymer, is not affected by the inorganic phase generated by the hydrolysis of the alkoxide. Furthermore, the excellent thermal stability of polyimides makes possible the heat treatment of the composite up to 350  C, allowing the inorganic network to increase its connectivity. Also in this case, however, the obtainment of nanometer and well-dispersed particles is not trivial. The interaction between the inorganic and the organic phases must be pursued through specific strategies: using polyamic acid containing pendent hydroxyl groups (Huang and Gu 2003) or siloxane functionalities (Park et al. 2002), which partially crosslink with the inorganic network, or using γ-glycidoxypropyltrimethoxysilane (GLYMO) as a compatibilizer to enhance the interchain interaction (Musto et al. 2004; Kizilkaya et al. 2010). Figure 8 shows the difference in the microstructure of the nanocomposite adopting or not a compatibilizer. Li et al. (2007) demonstrated that also in the case of PI/silica nanocomposites the order by which the chemicals are introduced in the batch may have a large effect on the size of the particles in the matrix. The better compatibility and interfacial strength between silica nanoparticles and the polyimide matrix bring to better mechanical properties, as demonstrated by Chen et al. (2004). The best performances were obtained with 10–15 wt% silica and the use of GLYMO, which allowed the formation of chemical bonds between the hydroxyl groups of the polymer and the epoxy groups of GLYMO. The maximum tensile strength increased from 65 to 90 MPa, and the maximum elongation at break increased from 7% to 10%. For a review on this subject, see Ragosta and Musto (2009).

Fig. 8 Different microstructures obtained in polyimide/SiO2 nanocomposites by using (right) or not (left) GPTMS as a compatibilizer (Polymer, 45, Musto P, Ragosta G, Scarinzi G, Mascia L, Polyimide-silica nanocomposites: spectroscopic, morphological and mechanical investigations, 1697–1706, 2004, with permission from Elsevier)

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An interesting example of interaction between the reactions leading to the organic and the inorganic polymerization, and the possible strategy to avoid it, is offered by nanocomposites based on polyurethane hybrid matrices and was described by Becker-Willinger (2014). Isocyanate polymerization may be affected by side reactions with small nucleophilic molecules such as water and amines, resulting in less cross-linking and affecting the mechanical properties in the final material. This effect may be reduced by blocking the electrophilic reactive functionalities of isocyanate molecules with molecules having a moderate nucleophilic character, thus protecting them from undesired side reactions. The bond between the isocyanate and the blocking molecule has to be strong enough to stabilize the sol before its application, but it must be cleavable during the final curing of the material. Further examples are rubbers reinforced with silica nanoparticles (Ykeda et al. 1997, Ykeda and Kameda 2004; Kohjiya and Ykeda 2003). Raw rubber or rubber vulcanizate may be swelled in tetraethoxysilane (TEOS), which is then transformed into silica by sol-gel reactions. In this case, the dispersed phase is synthesized within the matrix. A homogeneous distribution of nanoparticles is obtained, and mechanical performances are better with respect to the same nanocomposite obtained by mechanical mixing.

Applications of Sol-Gel Nanocomposites The development of sol-gel nanocomposite materials has been mainly driven by their potential use in many applications, and commercial sol-gel nanocomposites are already part of manufacturing technology. It is possible to define a commercial product as the one that is both offered for sale and used in the regular production of a device or item in general commerce, as already pointed out by B. Arkles (Arkles 2001). In this chapter, we present examples of sol-gel nanocomposites that potentially can have commercial application, based on the literature of the last 10 years.

Biomaterials Sol-gel nanocomposites are widely used as biomaterial (e.g., for drug delivery systems, cell encapsulating matrices, and artificial bone substitutes) because of their controllable physicochemical properties including size, morphology, composition, and microstructure and also for their controllable biocompatibility and biodegradability. A biomaterial can be defined as any matter, surface, or construct that interacts with living systems. Bioactive glasses were first developed by Hench and co-workers in 1969 (Hench et al. 1971) and represent a group of surface reactive materials which are able to bond to bone in physiological environment (Boccaccini et al. 2010). Bioactive glasses most widely used in biomedical applications consist of a silicate network incorporating sodium, calcium, and phosphorus in different relative proportions. For

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example, 45S5 bioactive glass universally known as Bioglass ® has a composition of 45% SiO2, 24.5% Na2O, 24.5% CaO, and 6% P2O5 in wt.%. To provide additional functionality, glass and glass-ceramic matrices are combined with synthetic and natural polymers such as poly(L-lactic acid) (PLA), poly (lactic-co-glycolic acid) (PLGA), and polycaprolactone (PCL). Peter et al. (2010a) fabricated bioactive glass-ceramic nanoparticles via the sol-gel process, which were blended with chitosan and gelatin to obtain composite scaffolds having macropores of 150–300 μm in diameter. The obtained scaffolds showed better degradation and swelling properties, increased protein adsorption, and very good cell attachment and spreading, with increasing the bioactive nanoparticles content. The scaffolds were proposed as potential candidates for alveolar bone regeneration applications. Xu et al. (2011) synthesized porous glass-collagen-phosphatidylserine (BG-COL-PS) nanocomposite scaffolds via the freeze-drying technique. The scaffolds had interconnected porous structures with pore diameter up to about 300 μm and with 75% in porosity. According to the culturing rat mesenchymal stem cells, the presence of PS stimulated osteogenic differentiation. Boccaccini et al. (2005) fabricated the tubular foam scaffolds of PLGA/ Bioglass ® composite for tissue engineering scaffolds for the regeneration of tissues requiring a tubular shape scaffold, such as intestine, trachea, and blood vessels. Maquet et al. (2004) fabricated highly porous scaffolds based on PLGA obtaining polymer foams containing 45S5 Bioglass ® by freeze-drying procedure. The presence of the filler retarded the degradation of the scaffolds while it stimulated apatite deposition. Hong et al. (2009) have investigated a new family of composites combining PLA and sol-gel-derived bioactive glass-ceramic (BGC) nanoparticles. 3D porous scaffolds were prepared by thermally induced phase-separation combining PLA and different concentrations of BGC nanoparticles. The in vitro studies showed that composites containing BGC nanoparticles with lower phosphorous and higher silicon content have better bioactivity than those with lower silicon and higher phosphorous content. More recently, El-Kady et al. (2010) have developed sol-gelderived bioactive glass nanoparticles/PLA composites by using solid–liquid phaseseparation method combined with solvent extraction. They used a modified alkalimediated sol-gel route to obtain bioactive glass nanoparticles. The modified sol-gel method resulted in reduction of the gelation time to about a minute rather than days as in the traditional sol-gel process. Furthermore, fast gelation prevented the aggregation and growth of colloidal particles to sizes larger than 100 nm. The proposed method is thus capable of delivering nanoparticles of sizes less than 100 nm with minimum agglomeration. It was reported that the scaffold’s pore size decreased with the increase in the glass nanoparticles content. The in vitro studies revealed that the addition of bioactive glass nanoparticles improved the bioactivity of the scaffolds. Natural-based materials such as polysaccharides (starch, chitin, chitosan) or proteins (silk, collagen) can be used as polymer matrices to prepare sol-gel nanocomposites. Peter et al. (2010b) have synthesized scaffolds made of SiO2CaO-P2O5 bioactive glass-ceramic nanoparticle and chitosan by using lyophilization technique. The developed composite scaffolds demonstrated adequate swelling and

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degradation with the addition of the nanoparticles. In vitro studies showed the deposition of apatite on the surface of the composite scaffolds, indicating the bioactive nature of the composite scaffolds. The investigation of the in vitro behavior considering osteoblast-like cells (MG-63) indicated that cells attached on the pore walls of the scaffolds and showed initial signs of spreading. A series of composites of various morphologies, such as fibrous membranes and 3D porous scaffolds, were developed by compounding polymers and bioactive glass nanofiber. Kim et al. (2006a) were the first to develop a composite of PLA filled with sol-gel-derived bioactive glass as a nanoscale composite fiber by means of electrospinning. Nanocomposites with a dense nanofibrous network were achieved. It was observed that glass nanofibers were uniformly dispersed in the PLA matrix. The excellent bioactivity of the nanocomposite was confirmed in vitro within a simulated body fluid by the rapid induction of bonelike minerals onto the nanofiber surface. Kim et al. (2006b) also developed bioactive glass nanofiber collagen nanocomposite both in the form of a thin membrane and as macroporous scaffold.

Chemical Sensors Sol-gel nanocomposites are widely used as sensitive material in chemical sensors. Sol-gel is relatively simple to transfer to an industrial scale and is particularly suitable for the deposition of thin films, as it is based on the deposition of a thin layer of the liquid precursor solution over a substrate at room temperature (by dip-coating, spin-coating, spraying techniques), followed by a thermal treatment. In the case of thin films, where the cost of precursors is minimized, the method is also not expensive. For these reasons, sol-gel may offer interesting challenges in the field of chemical sensors, and indeed it is actually largely investigated. An example of the potentiality of the sol-gel method is offered by the synthesis of nanocomposite thin films constituted of a nanoporous amorphous silica matrix and homogeneously distributed semiconducting metal oxide and/or metal nanoparticles. The nanoporosity provides the path for gas molecules to reach all the volume of the sensing material and the possibility for the nanoparticles to be efficiently exposed to the analyte. Based on this strategy, Guo and Tao (2007) developed ammonia optical sensors by coating an optical fiber with porous silica film containing Ag nanoparticles. In this case, the interaction of the NH3 gas induces a variation of the surface plasmon resonance peak of the Ag nanoparticles. Hasani et al. (2010) dispersed highly luminescent CdSe@ZnS core-shell quantum dots in hybrid organic–inorganic sol-gel matrix obtained from tetramethoxysilane and 3-aminopropyltrietoxysilane. The quenching of the luminescence of fragment of the nanocomposite xerogel has been monitored in the presence of different volatile organic compounds. Martucci at al. were the first to develop optical gas sensors based on metal oxide nanoparticles (NiO or Co3O4) dispersed in a porous SiO2 solgel film (Martucci et al. 2004a, b). The optical transmittance of the nanocomposite films increases upon exposure to CO or H2 gas due to the interaction of the gas

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molecules with the adsorbed oxygen on the metal oxide nanoparticles surfaces. The sensor sensitivity and selectivity have been improved by the addition of Au nanoparticles, which catalyze the reaction with the target gases and modified the absorbance spectrum. In fact the presence of the Au surface plasmon resonance peak permits to selectively recognize H2 and CO due to the wavelength dependence of the gas-induced optical transmittance variation (Buso et al. 2007b). Moreover by tailoring the thermal annealing conditions, it was possible to synthesize cookie-like Au/NiO nanoparticles showing unique structural and optical properties (Mattei et al. 2007). The high loading of metal oxide nanoparticles in the SiO2 matrix (up to 40 mol%) obtained by Martucci et al. (Cantalini et al. 2005; Della Gaspera et al. 2010b) allowed the realization also of conductometric gas sensors. They synthesized NiO, Co3O4, and SnO2 nanocomposite films which have shown a p-type (NiO and Co3O4) and an n-type response (SnO2) with greater sensitivity to H2 gas than CO and a detection limit of 10 ppm. Buso et al. (2007b) developed TiO2 sol-gel films containing Au nanoparticles for CO and H2 optical and electrical gas sensors. The authors present outstanding dynamics for hydrogen detection and a CO response dependent on the amorphous or crystalline (anatase) phase of the TiO2 film. Recently Lu et al. (2011) synthesized SiO2-graphene oxide nanocomposite decorated with Ag nanoparticles that can be used as biosensor for the glucose detection in human blood serum.

Catalysis Catalysis is the increase in the rate of a chemical reaction due to the participation of a substance called a catalyst. Unlike other reagents in the chemical reaction, a catalyst is not consumed. With a catalyst, less free energy is required to reach the transition state, but the total free energy from reactants to products does not change. The production of most industrially important chemicals involves catalysis. Similarly, most biochemically significant processes are catalyzed. For example, Pagliaro et al. (2011) synthesized sol-gel catalyst made of nanostructured Pd organosilica xerogels suitable for the catalysis of carbon–carbon bond formation in relevant organic reactions. Such nanocomposite has been commercialized with the trademarked SiliaCat Pd0 Hydrogel. Recently Sabada et al. (2014) developed a nanocomposite catalyst for the conversion of xylose to furfural based on aminopropyl-functionalized SiO2 sol-gel nanoparticles embedded in poly (styrenesulfonic acid). Photocatalyst is a catalyst that participates and modifies the reaction rate of chemical reactions under light irradiation. Since the discovery that TiO2 particles were capable of trapping and oxidizing organic compounds to minerals and small molecules such as CO2, extensive studies have been conducted to explore the potential of TiO2 as photocatalyst for environmental decontamination such as water purification and air cleaning. Nanocomposite based on TiO2 has been studied extensively because they can have better performances than pure TiO2.

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Li et al. (2010) synthesized mesoporous silica-supported Cu/TiO2 nanocomposites through a one-pot sol-gel method for the photoreduction of CO2 to hydrocarbons. The high surface area mesoporous silica substrate (>300 m2/g) greatly enhanced CO2 photoreduction, possibly due to improved TiO2 dispersion and increased adsorption of CO2 and H2O on the catalyst. Carbon nanotube/TiO2 nanocomposites were fabricated by Gao et al. (2008) by a surfactant wrapping sol-gel method that led to a uniform and well-defined nanometer-scale titania layer on individual carbon nanotube. The nanocomposites were found to enhance the initial oxidation rate of methylene blue by onefold compared to the pure TiO2 sample. This larger degree of rate enhancement is attributed to the supporting role of the carbon nanotubes and surface properties. Ag-TiO2 nanocomposites have been used as bactericide material, and recently it was demonstrated a long-lasting antibacterial activities of Ag-TiO2 nanocomposite sol-gel thin films under solar light irradiation (Akhavan 2009).

Wettability The wetting properties of a surface can be modified with appropriate surface coating which can impair desired wettability by tailoring the surface energy: lowering the surface energy, it will lower the wetting of the surface. Hydrophobic surfaces are those for which water is not wetting (spreading) the surface. Materials with very low surface energy are, for example, polydimethylsiloxanes (silicone) and polytetrafluorethylene (PTFE, Teflon®) with surface energies below 18 mN/m. Xu et al. (2010) developed a superhydrophobic (water contact angle larger than 150 ) highly transparent and stable organic–inorganic nanocomposite coating by a simple sol-gel dip-coating method. This method involves control of the aggregation of sol-gel SiO2 particles by polymerization and ultrasonic vibration to create a nanocomposite coating made of silica nanoparticles functionalized with aminopropyltriethoxysilane. Superhydrophobicity and transparency of the coating can be controlled by adjusting the initial concentration of monomer and the size of aggregates in the sol-gel. Thus, superhydrophobicity and high transparency can be concurrently achieved in a single coating. Lakshmi et al. (2012) studied the effect of the size of silica nanoparticles on the wetting properties of sol-gel nanocomposites obtained from methyltriethoxysilane. The coatings became superhydrophobic at an optimum silica concentration. The water repellence was further improved by the addition of fluoroalkylsilane (FAS). The optimum silica concentration was found to depend on the size of silica nanoparticles and FAS content. The superhydrophobic property of the coatings was due to the synergistic effect of surface chemistry and surface microstructure. Hao et al. (2012) used sol-gel nanocomposites made of polyhydromethylsiloxane and silica nanoparticles for developing superhydrophobic cotton fabric with favorable washing durability, maintaining its air permeability, color, and softness.

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Gao et al. (2014) developed a separation thin membrane made of single-walled carbon nanotube coated with sol-gel TiO2 that has superhydrophilic and underwater superoleophobic properties after UV-light irradiation. The robust and flexible nanocomposite films with a thickness and pore size of tens of nanometers can separate both surfactant-free and surfactant-stabilized oil-in-water emulsions in an ultrafast manner.

Scratch-Resistant Coatings Hard coating is one of the most studied applications for sol-gel nanocomposites coating, and it is one of the most promising for industrial application as pointed out by Sepeur and Frezer (2014). They reported that for a good abrasion resistance, the coating thickness of a cured sol-gel layer must be in the range between 3 and 20 μm. In fact, below 3 μm no significant increase in scratch resistance is observed, because the influence of the substrate. At a coating thickness of more than 20 μm, the inflexibility of the sol-gel coatings leads to cracks and poor adhesion on the substrate. In general, hard coating is based on SiO2, Al2O3, ZrO2, or TiO2 as network formers or fillers in the form of nanoparticles and organic-modified silanes such as epoxysilanes, methylsilanes, methacrylsilanes, and vinylsilanes. For example, Yuan et al. (2005) studied the influence of the size of SiO2 nanoparticles functionalized with 3-glycidoxypropyltrimethoxysilicane on the hardness of the nanocomposite coating deposited on aluminum substrate. More recently, Sowntharya et al. (2012) developed hybrid nanocomposite coatings from titanium tetraisopropoxide and epoxy- or acrylic-modified silanes for improving the resistance to scratch of polycarbonate substrate. The scratch as well as abrasion resistance increases as a function of coating thickness. The pencil scratch hardness improves from 2B for the bare PC to a maximum of 3H for the coating obtained from an aged sol derived from epoxy modified silane.

Conclusion Nanocomposites can be prepared by the sol-gel method in different ways and with different strategies. The general principles for their synthesis have been shortly discussed considering “in situ” and “ex situ” methods and different types of nanocomposites, with the matrix, the dispersed phase(s), or both obtained from sol-gel. The synthesis methods are discussed for nanocomposites with an oxide, with a hybrid organic–inorganic or a polymeric matrix, and with oxide, metal, or non-oxide nanoparticles. Representative examples of nanocomposites that potentially can have commercial application are described in the second part of the chapter, based on the literature of the last 10 years.

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Laura Brigo, Gioia Della Giustina, and Giovanna Brusatin

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Strategies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polymerizable/Cross-Linkable Hybrids . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Chelating/Complexing Agents . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Degradation of Organic Modification of Hybrids . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Organically Modified Dense Patterns . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Porous Patterns . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Inorganic Patterns . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Microfabrication and Nanofabrication . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Advanced Double-Tone Hybrid Materials Engineered for Multiple Pattern Technologies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

Hybrid organic–inorganic (HOI) sol-gel materials are becoming increasingly important as functional systems alternative to traditional organic resists for micro- and nanofabrication. Combining the processability of polymeric resists, with glass-like surface chemistry and enhanced mechanical/thermal stability, the direct use of hybrids as final micro-/nanostructured device materials is possible. L. Brigo Department of Industrial Engineering, University of Padova, Padova, Italy Center for Materials and Microsystems, Bruno Kessler Foundation, Trento, Italy e-mail: [email protected] G. Della Giustina · G. Brusatin (*) Department of Industrial Engineering, University of Padova, Padova, Italy e-mail: [email protected]; [email protected] # Springer International Publishing AG, part of Springer Nature 2018 L. Klein et al. (eds.), Handbook of Sol-Gel Science and Technology, https://doi.org/10.1007/978-3-319-32101-1_110

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Provided a specific engineering in material formulation, hybrid sol-gels can proficiently interact with radiation or undergo modifications under given thermal and pressure conditions. In this chapter, in addition to possible strategies to engineer hybrid sol-gel resists, an extensive presentation of developed material compositions and relative properties (mainly organic, inorganic, mesoporous) will be provided. Significant applications to traditional and next-generation lithography techniques will be shown, on the basis of recent reports from the scientific literature. Finally, few distinctive examples of devices realized through a synergistic engineering of hybrid materials and lithographic processes will be provided.

Introduction Hybrid organic–inorganic (HOI) sol-gel systems are an emerging class of resist materials for micro- and nanolithography. In the last two decades, several micro- and nanofabrication technologies have been used to pattern organic polymers or hybrid systems alternative to traditional photoresists. Photoresists have always and commonly been used to transfer patterns to inorganic materials for optics or electronics, the word “photoresist” originally indicating a photosensitive material that is at the same time resistant to acid etching after patterning. In fact, the resist is generally used as a sacrificial material deposited on the functional organic or inorganic material and patterned. The image of the sacrificial layer is then transferred to the functional material, in a pattern transfer step which more often uses radiation. This multistep process is time-consuming and makes the process complicated, often causing a deterioration of lithographic performance. Until recent days there was a limited availability of commercial polymeric photoresists, which properties are often limited by the temperature resistance, etching rate, and composition and confine their use to pattern transfer processes. With the aim to broaden the spectrum of spin-on materials that behave as resists but can also act as final device materials, HOI materials have become more and more important for micro- and nanofabrication technologies. HOI are built up by the bottom-up sol-gel approach at low process temperatures and offer a broad range of compositions and structures, as outlined in Fig. 1. First, they proficiently interact with radiation or undergo modifications under specific thermal and pressure conditions thanks to the presence of organic modifications which can basically work as degradable or polymerizable/cross-linkable components, as described in section “Strategies”. Essential requisite, they are further processable with development steps after radiation interaction, as the chemical changes allow a different response of the irradiated and unirradiated areas against a proper developer used to chemically etch one of the two areas. This process allows structuring HOI surfaces by a direct lithographic process. Section “Materials” reports examples of resist compositions which can spun from mainly organic slightly inorganically modified to totally inorganic (mainly ceramic or glass-like) in which the organic modification has mainly the role of enabling photosensitivity or molding processes.

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Fig. 1 In the periodic table metals (Me) whose alkoxides are commercially available are indicated in yellow. They are processed by a bottom-up sol-gel chemistry, together with organic components or organically modified silicon alkoxides, and directly patterned by different lithographic tools

From the point of view of micro- and nanofabrication, they combine polymer-like fabrication processes, typical of polymeric resists, with glass-like surface chemistry that is beneficial for many applications and allow their direct use as the final patterned material. Literature on fabrication with different lithographic tools used to direct pattern HOI sol-gel films, during the last years, is rather extensive, as can be seen from examples reported in the first two sections. These techniques use electromagnetic radiation, from the most common UV light (for laser writing or mask-assisted UV lithography) to X-rays and electrons or preformed stamps to generate patterns by nanoimprint lithography (NIL). However, alternative micro- and nanolithographies have emerged in the last years. Among the so-called next-generation lithographies, extreme ultraviolet (EUV) lithography is the most promising as high-resolution high-throughput technology for advanced silicon chip technology fabrication. 3D fabrication techniques are attracting more and more attention, and in particular, two-photon submicron fabrication of hybrid sol-gel materials has recently appeared in the scientific literature and has already demonstrated its powerful potentialities. Some distinctive examples of HOI resist lithography based on the advanced use of the lithographic techniques combined to engineered hybrid materials are reported in section “Microfabrication and Nanofabrication.” In addition, nowadays, much effort is put on the scaling up of currently developed lithographic tools for industrialization and direct application to device fabrication. Roll-to-roll (R2R) and EUV techniques

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represent emblematic examples, and some examples of hybrid sol-gel materials applied to these technologies have been reported. Our presentation of hybrid materials developed for micro- and nanofabrication technologies concludes, in section “Applications,” with an on-purpose non-exhaustive report of a few examples of devices realized through a synergistic engineering of HOI materials and lithographic processes.

Strategies Interaction between radiation and material and material rheology are key points for a good lithographic outcome. The general rule to design a new resist material is to select a formulation having network functionalities sensitive to the source (photons, electrons, heat) of the lithographic technique. The goal is to create a chemical change (organic or inorganic cross-linking, organic degradation, hydrophobicity/ hydrophilicity) in the material structure. Especially in radiation and electron-assisted lithography, this modification produces a contrast between the exposed and not exposed areas, exploitable to realize patterns on films by a wet etching process (developing step). In replication technologies, such as thermal and UV-assisted NIL, a further constraint has to be satisfied during materials development. Rheology and system viscosity have to be carefully tailored to achieve good quality and high replication fidelity. The main adopted strategies can be grouped into the development of hybrid materials having: – Cross-linkable or polymerizable units – Chelating/complexing agents – Degradable moieties or molecules

Polymerizable/Cross-Linkable Hybrids This strategy is mainly based on the use of organically modified silicon alkoxides, easily available in the market, whose Si-C bond is stable upon hydrolysis during solgel process. In particular, organofunctional silanes are compounds combining the reactivity of an organic group and the inorganic functionality in a single molecule and covalently bonding the inorganic and organic components. The formation of reactive silanols in the presence of moisture leads to the development of a siloxane network bearing pendant organic functional units. By careful control of the synthesis conditions, storage-stable resins with negative resist behavior can be synthesized for a large variety of applications and particularly to develop also commercial materials. Several examples of direct patterning of vinyl-, acrylate-, or methacrylate-based hybrids are reported by using different lithographic techniques (Oubaha et al. 2006; Moreira et al. 2005).

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'R

R' Si

HO

R Si

'R

O MeO

R

R' Si

OSi MeO

Si O

MeOH

R

in situ H2O

Si MeO

R = 3-methacryloxypropyl R' = phenyl

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R'

'R Si

OSi MeO

OH MeOH

R Si

O

R

T3 OSi O

Si OMe OMe

OMe OMe

Fig. 2 Proposed mechanism for establishing T3 species (Buestrich et al. 2001)

ORMOCER® materials (organically modified ceramics developed by Fraunhofer Institute) are one of the most famous examples of class II dense HOI systems, largely used in micro- and nanofabrication. They are synthesized by hydrolysis and polycondensation reactions of alkoxysilanes with organic cross-linkable substituents, as shown in Fig. 2, and are UV (broadband + i-line) or thermally curable (~120
C). Thanks to the presence of polymerizable double bond C=C functionalities (Fig. 2), patterning of thick (10–50 μm) hybrid structures by UV lithography, hot embossing, two-photon lithography, and NIL is reported (Buestrich et al. 2001; Ovsianikov et al. 2007). The addition of a proper initiator is also necessary to promote the organic curing. Free radical photoinitiators – often belonging to the commercial family of Irgacure products (BASF) – are largely used with unsaturated hybrid systems (e.g., acrylic and methacrylic units). Examples of commercial photosensitive resists developed exploiting this approach are OrmoComp, a proprietary UV-curable acrylate sol-gel hybrid resin from MRT (micro resist technology GmbH), suitable for the fabrication of microoptical components and photonics applications. Other products from MRT, such as OrmoStamp and OrmoCore, belong to the same class of hybrid polymers and are used for production of waveguides or NIL stamps at wafer level, respectively. If photosensitive unsaturated HOI materials already reached the market, systems based on epoxy modified silicon alkoxides are still under development. Patterning of epoxy hybrid sol-gel materials has been largely studied by several authors (Brusatin et al. 2006; Della Giustina et al. 2007; Pina-Hernandez et al. 2010; Dal Zilio et al. 2010), and the most common sol-gel precursors are 3-glycidoxypropyltrimethoxysilane (GPTMS) and 2-3,4-epoxycyclohexylethyltrimethoxysilane (EETMOS). The developed strategies allowed to obtain hybrid SiO2-epoxy-based dense materials (Fig. 3), with good optical properties and patternable with different fabrication technologies, such as UV lithography, XRL, and EBL (Della Giustina et al. 2008; Jabbour et al. 2008). A key issue of epoxy hybrid materials is the photocuring process. Epoxy ring photopolymerization occurs by a cationic mechanism (Fig. 4) which is known to be insensitive to oxygen (Crivello and Ortiz 2001). This allows to overcome the inhibition problems and uncompleted polymerization of many acrylic systems (Versace et al. 2010).

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Fig. 3 Sol-gel reactions of organosilane and a transition metal precursors (GPTMS and tetraethyl orthogermanate, TEOG) with the formation of an organically modified oxide network. Development of oligomeric polyether chain after UV curing of epoxy rings

Fig. 4 Polymerization of epoxy ring in a cationic way (Jabbour et al. 2008)

Common cationic photoinitiators are photoacid generator (PAG) molecules based on iodonium or triarylsulfonium salts (Crivello 1984). The photolysis mechanism produces a protonic acid which propagates the reactions in chain by nucleophilic attack to epoxy monomer and leads to a formation of the organic network. Recently nonionic molecules have been also studied as PAG initiators for cationic polymerization of 3-glycidoxypropyltrimethoxysilane HOI sol-gel materials (Torti et al. 2015). However, some nonionic commercial initiators for lithographic applications are starting to appear even in the market (e.g., the iminotosylate Irgacure PAG 121-BASF).

Chelating/Complexing Agents A versatile scheme to pattern HOI materials is the use of chelating or complexation agents (Dinachali et al. 2013). This approach is mainly exploited to prepare hybrid materials starting from transition metal alkoxides precursors in order to obtain organically modified metal oxide different from silicon one.

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Sensitive metal-organic complexes have been often formed by the reaction of metal alkoxides with organic molecules such as β-diketone, β-ketoester, or carboxylic acids usually in an alcohol medium. In fact, metal alkoxides are very reactive compounds due to the presence of electronegative alkoxy groups that cause the metal atom to be highly prone to nucleophilic attack. These molecules act as a Lewis base that donates a pair of electrons (electron-pair donors) to the transition metal via a coordinate covalent bond or as ions which surround the metal ions in a complex: thanks to the empty valence orbitals, transition metal ions can behave as Lewis acids (electron-pair acceptors). In this way, the addition of these chelating agents slows down hydrolysis and condensation reactions due to steric hindrance and imparts to the system the ability to be patterned by numerous techniques such as UV lithography, EBL, and thermal or UV-assisted NIL (Garoli and Della Giustina 2015; Saifullah et al. 2010; Dinachali et al. 2013) forming a new photosensitive metalorganic precursor. Methacrylic acid (MAA), 2-(methacryloyloxy)ethyl acetoacetate (MAEAA), benzoylacetone (BzAc), ethyl acetoacetate (EAcAc), and hydroxyl-substituted aromatic ketones are some examples of chelating agents used in literature (Saifullah et al. 2003; Tohge and Takama 1999). Figure 5 reports an example of Ta modification by MAEAA, but the same strategy can be used for many different oxides. Al2O3, Ga2O3, In2O3, Y2O3, B2O3, TiO2, SnO2, ZrO2, GeO2, HfO2, Nb2O5, Ta2O5, V2O5, and WO3 are various oxides already patterned by exploiting this synthetic platform (Dinachali et al. 2013). The MAEAA β-ketoester is a bifunctional chelating agent bearing a methacrylate group which provides the reactive group for a free radical polymerization. On the other side, the carbonyl C=O groups form a stable metal complex with a bidentate character, as confirmed by the FT-IR spectra of Fig. 5. The obtained metalorganic precursor is a stable, low viscosity solution and allows to overcome shorter shelf life problems typical of some other methacrylate-modified systems. The liquid metalorganic precursors obtained by chelated monomer route provide the proper rheology and controlled viscosity with no need of additional solvent for replication technologies. Taking advantages from these properties, different oxide nanostructures have been fabricated by thermal and step-and-flash imprint lithography (Dinachali et al. 2013). Another example of chemical modification of metal alkoxides to obtain photosensitive material is described in Fig. 6. Vinylpyrrolidone is used to slow down the condensation of hydrolyzed Ti(O-nBu)4 and to form a complex by an H-bonding interaction of carbonyl C=O groups with Ti–OH. This leads to the formation of acrylate-modified TiO2 network cross-linkable by the UV light (Segawa et al. 2004b). UV irradiation causes both the polymerization of C=C bond and the weakening of hydrogen bond of C=O groups. The synthetic approach underlying the abovementioned examples is the same; however, different phenomena occur according to the chelating molecules and the used lithographic techniques. In literature, two principal mechanisms of pattern oxide formation can be highlighted:

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– Organic polymerization of unsaturated groups, in case of ligands bearing polymerizable moieties and forming stable complexes, as in the case of MAEAA and vinylpyrrolidone. An organic polymer can be established by using proper free radical initiators and by the application of temperature or radiation sources (e.g., thermal or UV-assisted NIL).

Fig. 5 (continued)

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– Photolysis and complex breaking in case of ligands making weak complexes, as in the case of MAA or BzAc. Under EB or UV radiation, chelated rings undergo a photodissociation bringing to an inorganic cross-linking, as described in Fig. 7 (Stehlin et al. 2014; Saifullah et al. 2010). This effect is even more pronounced in chelating/complexing molecules having an absorption band in the irradiation range. For example, photosensitive ZrO2 sol-gel films modified with ethyl acetoacetate, acetylacetone, benzoylacetone,

Fig. 5 (a) Clear, stable, and slightly colored metal oxide precursors formed when alcohol by-product was removed following the reaction shown in (c)-(3). (b) Characteristic infrared absorption peaks of metal oxide precursors formed by reacting metal alkoxide and MAEAA in a 1:2 ratio. The broad vibration bands corresponding to particular bonds are indicated on top. (c) The β-ketoester group undergoes keto-enol tautomerism (1); the enol form of MAEAA is stabilized by chelation with tantalum ethoxide that results in the formation of a chemically stable chelated alkoxide: stoichiometric replacement of an alkoxy group by a β-ketoester ligand (2). The general reaction of an alkoxide with MAEAA is also shown (3) (Dinachali et al. 2013)

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Fig. 6 Schematic illustration of capping of the Ti alkoxide by VP in the gel film before (a) and (b) after UV irradiation (Segawa et al. 2004)

Fig. 7 Schematic view of the DUV material preparation, from MOC thin films to a metal oxide thin film, with proposition of a potential cross-linking molecular scheme and complete mineralization by DUV irradiation in the case of TiOC (Stehlin et al. 2014)

dibenzoylmethane (DBzMe), and 1-hydroxy-2-acetonaphthone (HAN) show absorption maxima at 280, 302, 334, 365, and 410 nm, respectively (Noma et al. 2004). Therefore, the material absorption range can be controlled selecting the proper chemical structure of chelating agent. Thanks to the elimination of the complexing molecules by photodissociation of chelated rings and development of

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volatile products, inorganic cross-linking proceeds by the hydrolysis and condensation reactions of metal oxide precursor (Saifullah et al. 1999) or other cross-linking mechanisms involving radical species (Stehlin et al. 2014).

Degradation of Organic Modification of Hybrids Unlike the first two strategies where the main mechanism is organic polymerization or inorganic cross-linking, the last approach exploits the degradation of organic moieties or molecules to produce a modification in material structure, as required in micro- and nanofabrication technologies, and a final inorganic pattern is usually achieved. In fact, extending the ability of directly micro- and nanostructuring also metal oxides is of great interest in many applications such as electronic, optoelectronic, and photonic devices. HOI sol-gel materials prepared from phenylsilane such as diphenylsilanediol, phenyltrimethoxysilane, and 4-bis(triethoxysilyl)benzene are one of the classes of photodegradable resists (Brigo et al. 2011b; Falcaro et al. 2011). Due to ρ-π conjugation between the silicon atom and the phenyl substituent, these compounds possess an absorption band in the short UV wavelength range (around 280 nm), and their use allows to synthesize innovative photosensitive materials (Sato et al. 2003). An example of phenyl-modified silica network is a hybrid sol-gel material prepared from 1,4-bis(triethoxysilyl)benzene, a bridged polysilsesquioxane compound presenting an aryl bridge between two trifunctional silicon groups (Brigo et al. 2011b). Figure 8 shows as UV exposure causes a decrease of absorption bands at 220 nm and at 270 nm, characteristic of the phenyl presence. The general interaction mechanism is disconnection of the r-p conjugation between silicon and the phenyl group leading to a progressive removal of aromatic moieties with UV light. As a consequence, the exposed material became more inorganic increasing the hydrophilic surface character and its solubility in acidic aqueous media. X-ray and EB irradiation induce a similar phenomenon with the gradual elimination of aromatic rings (Falcaro et al. 2011; Brigo et al. 2010): the change in

Fig. 8 UV-vis spectra of 1,4-bis(triethoxysilyl)benzene-based films (a) and detail of the mode at 270 nm (b) for different UV exposure times (Brigo et al. 2011b)

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Fig. 9 (a) Grazing incidence FT-IR spectra of Brij56 templated nanocomposite silica film before (dotted line) and after (solid line) exposure to deep-UV light for 180 min (Dattelbaum et al. 2005). (b) Schematic of the preparation for patterned mesostructured/mesoporous silica thin films. (c) Optical microscope image showing contrast difference between mesostructured (purple) and mesoporous (yellow) regions (Dattelbaum et al. 2003)

chemical composition and in surface properties generates the contrast necessary for micro- and nanofabrication. This strategy is often applied to pattern mixed MO2SiO2 network (M=Al, Zr, Hf, etc.) and to obtain almost completely inorganic structures (Della Giustina et al. 2015; Grenci et al. 2015). Among the photosensitive HOI materials exploiting organic photodegradation, mesostructured systems should be included. Surfactant used to create the organic templating phase can play the role of photodegradable component (Fig. 9a). Short-wavelength UV light generates oxygen radicals and ozone that can oxidize and remove the organic surfactant (Dattelbaum et al. 2005; Innocenzi et al. 2008). Photochemical calcination generally used to obtain mesoporous systems can also produce different reactivity and wetting properties. Therefore, patterns of mesoporous regions in a mesostructured film can be achieved by spatially controlled UV exposure, as reported in Fig. 9b, c (Dattelbaum et al. 2003). A complete selective removal of the mesostructured phase can also be performed by using a proper solvent. A totally different strategy to fabricate inorganic structures is the photocatalytic degradation of organic species linked to Ti or Zn oxides, instead of the progressive elimination of the organic component due to its direct absorption. This allows in principle to transform materials with different organic modification into an inorganic oxide through the photocatalytic process occurring during UV pattern (Zanchetta and Della 2014). In fact, in photocatalysis, the absorption of a photon of energy greater than the bandgap of the semiconductor excites an electron–hole pairs, which generate free radicals able to oxidize organic species producing mineral salts, CO2 and H2O. The oxidation reaction is represented by:

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Fig. 10 FT-IR spectra of the RT sample before (red line) and after different UVexposure times with an LC5 lamp from a distance of 3.5 cm, with a corresponding incident intensity (considering the entire emission spectrum) of 450 mW/cm2. The spectra of samples treated at 500
C and 800
C are also reported (brown and dark green spectra, respectively) (Gardin et al. 2010) hνEg ðTiO2 Þ

organic þ O2 ƒƒƒƒƒƒ! CO2 þ H2 O þ mineral acids: Figure 10 shows the progressive decrease of the FT-IR absorption bands in 3000–2800 cm1 interval due to the stretching vibration of CH2 and CH3 bonds, confirming the decomposition of the organic species in an epoxy-based silica–titania system (Gardin et al. 2010). This effect can be used in both UV lithography and UV-assisted NIL, as described in section “Materials.”

Materials HOI materials guarantee solution processability with, at the same time, high lithographic performance and a wide choice of properties: among them, thermal resistance (typically up to 300
C), mechanical resistance, chemical endurance, and low dimensional shrinkage. Furthermore, functional properties such as optical, electrical, porosity, etc. can be tuned, and specific functions can be achieved by embedding nanoparticles, dyes, or other active molecules. Organic functionalities, such as epoxy, acrylate, phenyl, chelating, or coordinating agents, play the main role toward the lithographic tool, as described in the

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previous section, while metal alkoxides are generally selected for the focused applications, allowing the preparation of thermal-, pressure-, and radiation-sensitive resists in a wide range of ceramic compositions, from the most commonly found SiO2 to GeO2, TiO2, ZrO2, HfO2, Al2O3, PZT, etc. These different oxides are coupled with organic functionalities starting from small molecules and liquid precursors, allowing a high degree of control over composition, purity, structure, and processability. Common characteristic of HOI materials is that they possess a double polymerization ability: a cross-linked network is formed by photo- or thermal cross-linking of the organic part, and an inorganic network of strong covalent bonds is derived from hydrolysis and condensation of metal alkoxides. As a consequence, HOI resists often present a good pattern shape preservation after lithography as well as high replication fidelity in replica molding technologies. As often claimed, this also could determine low/negligible shrinkage during photocross-linking (Ovsianikov et al. 2008), but this consequence should be carefully considered. With this respect, few data are available, and shrinkage as low as 6% are reported (Pina-Hernandez et al. 2010). In the following, most significant examples of HOI resists used to achieve dense, porous, and inorganic patterns are discussed.

Organically Modified Dense Patterns First developed HOI resists compositions were mainly based on SiO2, synthesized from organically modified silicon alkoxides, with organic functions enabling photosensitivity or molding process through photo- or thermal polymerization. Often, a minor part of metal-organic precursors different from Si is added, in order to confer optical, mechanical, or other functional properties to the patterned films depending on their nature and concentration. Hybrid compositions based on organically modified silicon alkoxides allow to tailor the rheology of the film during the imprinting process (Gale et al. 2005; Peroz et al. 2009; Letailleur et al. 2010), while inorganic cross-linking and organic polymerization determine the hardening of the patterned structures, thermally or by radiation. In these systems the organic component remains in the final pattern, assuring the achievement of dense material structures. The silica network, together with a fully densified structure, accounts for the good optical quality of these glassy hybrid films that were largely used to fabricate channel optical waveguides and other optical systems such as gratings, lenses, etc. (Krug et al. 1992; Luo et al. 2005; Oubaha et al. 2006; Dal Zilio et al. 2010; Moreira et al. 2005). A typical example is the composition based on methacryloxypropyl trimethoxysilane (MAPTMS) and methacrylic acid (MA), both of which had photopolymerizable C=C double bond groups, and other metal oxides, typically Zr or Ti n-propoxide, as an inorganic network former. Several papers report on similar compositions (Vicente et al. 2014; Segawa et al. 2006; Kim et al. 2005b, Luo et al. 2005), in which the presence of TiO2 or ZrO2 determines refractive index modulation.

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Beside the presence of silica or other metal oxides, distinguished feature is the possibility of incorporating other functional species, as fluorinated groups (i.e., to control the refractive index, achieve electrical or magnetic properties, hydrophilicity, or hydrophobicity; etc.) (Moreira et al. 2005; Tadanaga et al. 2004; Garoli and Della Giustina 2015) and functional organic dyes or molecules, thanks to the low T synthesis, allowing to get directly patternable functional micro- and nanostructures (Brusatin et al. 2008; Brigo et al. 2012). As an example, hybrid in which fluorinated functionalities are covalently bound to the silica network were prepared (Kim et al. 2005a), as shown in Fig. 11, allowing the refractive index control. By varying their contents, refractive index of the synthesized resin is then controlled (1.469–1.513 at 850 nm). Simultaneously, minor volume contraction behavior during microscale molding is achieved thanks to the designed hybrid composition. A large core optical waveguide structure is formed by these hybrid materials, which exhibit low near-IR wavelength absorption, with measured optical propagation losses lower than 0.25 dB cm2. Among dense hybrid resist compositions, ORMOCER ®s, whose general chemical structure is shown in Fig. 12, have been largely developed and optimized (Buestrich et al. 2001; Streppel et al. 2003). Commercial OrmoComp ® and related hybrid resists offer glass-like material properties after UV curing and are standard products for a broad range of pattern sizes and component dimensions, for UV and

Fig. 11 (continued)

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O

F2 C O

Si(OMe)3

C F2

(MeO)3Si

F2 C

F2 C C F2

F2 C C CF3 F2

+ OH

HO Si

Catalyst & Temperature CF3 F2C

CF2

F2C

F2C

Si

O

CF2

O

CH3 O

CF2

O

F2C

Si O O

Si Si

O

O

Si

O

O

O

Si O

CF2

Si

F2C

CF2 F2C

CF2 F2C

CF2 CF3

O O Si O O

O

Fig. 11 (a) Design of the large core multimode optical waveguide, (b) SEM image of the cleaved edge of one waveguide, (c) SEM image of the fabricated ridge waveguide array, and (d) the light propagated image of the waveguide array. (e) Reaction scheme of silane precursors for the synthesis of fluorinated non-hydrolytic sol-gel hybrid materials (Kim et al. 2005a)

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Fig. 12 Chemical structure of ORMOCER ® polymers. Methoxy (Me) and phenyl (Ph) groups are linked to the inorganic Si-O backbone. R, organic side chain (Streppel et al. 2003)

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R Si O

O OMe

Ph

Ph Si

Si O n

Fig. 13 SEM image of the test geometries for defining the dynamic range of the photosensitive polymers with different photoinitiators. In each picture, the laser power was increased from right to left at 4 mW steps. (a) Irgacure OXE01 (12–36 mW), (b) Irgacure 819 (16–36 mW), (c) Darocur 4265 (16–36 mW) (Harnisch et al. 2015)

Fig. 14 SEM image of the prism array manufactured from OrmoComp ® (Harnisch et al. 2015)

UV-assisted molding, and recently ink-jet printing epoxy-based formulations (Voigt et al. 2011; Kim et al. 2011). Figure 13 shows an example of material optimization for two-photon polymerization (2PP) lithography (Harnisch et al. 2015); in particular, the influence of different photoinitiators on the 2PP behavior of OrmoComp ® is studied. A more complex structure in the form of a prism array was produced on glass (Fig. 14). The special features of this prism array are the individual inclination-and tilt-angles of each prism: very defined edges and overall smooth prism surfaces were obtained. HOI patternable materials were also proposed as inserts alternative to the metallic structures for the manufacturing of microstructured master molds in microinjection molding, a key technology for mass production of microstructured

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Fig. 15 Direct realization of microstructured stamps for microinjection molding through UV lithography: (a) spin coating of TMSPM-Zr solution on the metallic substrate and UV exposure of the film through a mask reporting the required structures; (b) final stamp obtained after development; (c) microinjection molding of PS substrates using the zirconia-based directly patterned stamp; (d) PS microstructured stamp obtained with microinjection molding; (e) MSC adhesion on the PS microstructured substrate; (f) MSC differentiation into osteoblasts on the PS microstructured substrate; (g) SEM image of the microstructured hybrid stamp; (h) the PS replica; and (i) cells adhered onto the PS substrate (Zanchetta et al. 2015)

surfaces. The use of directly patterned stamps presents a great advantage on the overall manufacturing process as it allows a fast and simple one-step process with respect to the use of milling, laser machining, electroforming techniques, or conventional lithographic processes for stamp fabrication. By using 3-(trimethoxysilyl)propyl methacrylate (TMSPM)-based dense HOI in combination with ZrO2 precursors and methacrylate monomers, hybrid films were UV patterned directly on steel substrates. In this study, different diameter micrometric holes were produced, then replicated in polystyrene in a short time through microinjection molding (Fig. 15) that were used to study the effects of the substrate topography on hMSCs adhesion, proliferation, and osteogenic differentiation (Lucchetta et al. 2015; Zanchetta et al. 2015). For these studies the

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Fig. 16 Synthesis of the polyimide–titania precursor and hybrid thin films (Chang et al. 2009)

availability of a number of identical substrates was a key point, to perform multiple biological analysis and verifying reproducibility. In case HOI resist does not contain organically modified silicon alkoxides, dense HOI patterned structures are also achieved using monomers, as methacrylic acid or vinylpyrrolidone, in combination with metal alkoxides (Segawa et al. 2004; Garoli and Della Giustina 2015; Segawa et al. 2004; Segawa et al. 2010), or modifying metal alkoxides with different strategies. An example of alternative synthesis, without using any coupling or chelating agents, is reported in Figs. 16 and 17. A polyimide–nanocrystalline titania hybrid high refractive index optical material is prepared, whose photosensitivity is provided by grafted methacrylate groups, exhibiting nanocrystallinity and patternability (Chang et al. 2009).

Porous Patterns Self-assembled mesoporous thin films are bottom-up synthesized materials for which lithographic processing allows to obtain hierarchically structured materials, as organization resides on multiple length scales: porosity scale, lithographic feature scale, film thickness scale, and fabricated area. Such complex systems that selfassemble into organized structures can be patterned by several strategies. Brinker et al. (Doshi et al. 2000) were the first to demonstrate the possibility of patterning mesoporous films by lithographic techniques, exploiting either mesophase change or mesostructure disruption in refractive index, pore size, surface

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Fig. 17 (a) SEM images on the patterned photosensitive polyimide–titania hybrid film; (b) HRTEM images of the studies hybrid films: circles indicate nanocrystalline titania in the hybrid film (Chang et al. 2009) Fig. 18 Irradiation of the PAG, a diaryliodonium salt, at 256 nm, results in photodecomposition to yield the Brønsted superacid, H+SbF6 plus an iodoaromatic compound, and organic by-products. Thus, UV exposure of the photosensitive mesophase through a mask creates patterned regions of differing acid concentrations compartmentalized within the silica mesophase (Doshi et al. 2000)

area, and wetting behavior occurring upon UV irradiation (Fig. 18). Hybrid silica sol containing a diblock copolymer as the surfactant and a photoacid generator (PAG) were synthesized. Upon UV curing, patterns were created by a mask on the silica films: the acid generation through UV light (Fig. 19) gives a higher silica polycondensation rate with respect to the part of the film that is not exposed. The cured silica network is more interconnected and more durable: the selective etching and the formation of structures with a preserved mesophase are allowed. Patterned structures can be achieved by changing not only the condensation state but also the mesophase. Changes in the mesophase are commonly observed in mesoporous silica films during the thermal treatment stage because of contraction in the direction normal to the substrate; however, the phenomenon can also be induced by UV light for the films with high surfactant concentration (Innocenzi et al. 2008).

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Fig. 19 UV patterning of mesostructured silica films. (a) Several fabrication pathways are possible: (b) after illumination with UV light, the unexposed part of the film is selectively wet etched; (c) the UV curing induces a mesostructure phase transformation in the exposed regions; the two regions have different refractive indexes; (d) the film is patterned with UV light and then thermally treated; this process produces a change in the refractive index between the irradiated and unexposed regions but not a change in the mesophase (Innocenzi et al. 2008)

Sophisticated lithographic techniques are adopted to achieve higher aspect ratios, higher resolutions, and well-defined complex patterns. For instance, deep X-ray lithography (DXRL) was demonstrated to allow a highly controlled patterning of mesoporous silica films (Falcaro et al. 2008). A simultaneous increase of silica polycondensation and partial removal of the templating agent were induced by synchrotron radiation, without any structural change in mesophase or damage. The areas of the film that were not exposed to radiation could be selectively etched owing to a lower cross-linking degree of the inorganic network. In the reported research work, a step forward toward the realization of functional microstructures was made employing a dip-pen technique, using the tip of an atomic force microscope and rhodamine 6G as ink, to obtain a specific functionalization on single mesoporous pillars fabricated by DXRL. Upon contact, capillary forces drove the solution from the cantilever tip into the pillar, filling mesopores with the rhodamine solution (Fig. 20).

Inorganic Patterns Highly inorganic (>80%) or completely inorganic final patterned structures can be micro- and nanofabricated starting from spin-on HOI resists with final thermal treatment, as in the illustrative example of PZT or PLZT patterns (Garoli and Della

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Fig. 20 (a) Confocal fluorescence image of patterned mesoporous pillars, where areas functionalized after patterning by dip-pen writing with rhodamine 6G appear in yellow; (b) for comparison, patterned mesoporous film fully impregnated by rhodamine 6G. (c) Optical profilometry of hollow pillars patterned by DXRL on mesoporous silica films (Falcaro et al. 2008)

2015; Weihua et al. 2003; Gaoyang et al. 2004; Tohge and Takama 1999; Bae et al. 2006; Park et al. 2010b; Benkler et al. 2014). However, strategies have been developed to achieve final inorganic compositions at low process temperatures, finding out direct, fast, economical, and easy ways to avoid the need for a multistep process and for standard photoresists. Main strategies use organically modified silicon alkoxides whose organic function acts as network modifier or as bridging group or metal alkoxide modified by chelating or coordinating species to patterning by radiation-assisted lithography and producing completely inorganic features after photodegradation of the organic component. Many of these examples can be found in literature, among them ZrO2, HfO2, TiO2, Al2O3 using mainly UV lithography, EBL, and NIL (Saifullah et al.

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2010; Dinachali et al. 2013; Ohya et al. 2002; Imao et al. 2006; Segawa et al. 2003; Park et al. 2010a; Zanchetta et al. 2013). A significant example is the direct pattern of inorganic ZnO nanocrystalline features, starting from a sol-gel hybrid composition of zinc acetate dehydrate, a chelating agent to stabilize the formation of metal oxide “clusters” (ethyl acetoacetate), and a tri-block copolymer to enhance the contrast between exposed and unexposed regions (Fig. 21) (Costacurta et al. 2011; Falcaro et al. 2009; Malfatti et al. 2015). Besides patterning, DXRL triggers the formation of a ZnO crystalline phase, which can be controlled by a thermal treatment at 200
C and is correlated with the presence of a residual organic compound in the matrix. Broad literature reports on patterning of inorganic TiO2, but only few paper reports on direct patterning of TiO2 films (see, e.g., Luo et al. 2005; Segawa et al. 2004; Segawa et al. 2003; Lim et al. 2010; Ganesan et al. 2012; Parashar et al. 2003). General strategies are based on the addition of organically modified silicon alkoxides and methacrylic monomers, to interact with radiation or control viscosity during imprinting or organic degradating agents such as chelate b-diketones (Barbè et al. 2012). However, a following calcination step is required to fully transform the features into inorganic structures, or a treatment at moderate temperatures in case of stabilized such as diethanolamine is used (Huang et al. 2016). Direct fabrication of completely inorganic titania or silica–titania micro-patterns is reported (Gardin et al. 2010; Gardin et al. 2011; Zanchetta et al. 2012; Zanchetta et al. 2014) starting from HOI spin-on titania resists. In this case, the small crystalline TiO2 clusters, shown in Fig. 22, produced during the sol-gel reaction generate oxygen radicals that can oxidize and decompose the surrounding organic compounds when exposed to UV. This photocatalytic effect converts the HOI in the exposed areas into inorganic TiO2 patterns after development (see Fig. 34) or after thermal imprinting, with refractive index up to 1.8 at 632 nm). An important class of hybrid resists was recently developed for extreme UV (EUV) lithography, one of the most important next-generation lithography for highvolume manufacturing; these HOI resists are based on small inorganic oxide clusters,