Advanced Lightweight Multifunctional Materials 0128185015, 9780128185018

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Woodhead Publishing in Materials

ADVANCED LIGHTWEIGHT MULTIFUNCTIONAL MATERIALS

Edited by

PEDRO COSTA CARLOS M COSTA SENENTXU LANCEROS-MENDEZ

Woodhead Publishing is an imprint of Elsevier The Officers’ Mess Business Centre, Royston Road, Duxford, CB22 4QH, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States The Boulevard, Langford Lane, Kidlington, OX5 1GB, United Kingdom Copyright Ó 2021 Elsevier Ltd. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library ISBN: 978-0-12-818501-8 For information on all Woodhead Publishing publications visit our website at https://www.elsevier.com/books-and-journals Publisher: Matthew Deans Acquisitions Editor: Gwen Jones Editorial Project Manager: Charlotte Rowley Production Project Manager: Surya Narayanan Jayachandran Cover Designer: Alan Studholme

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Contributors R. Alves Center of Physics, University of Minho, Campus de Gualtar, Braga, Portugal; Institute of Science and Innovation for Bio-Sustaninability (IB-S), University of Minho, Campus de Gualtar, Braga, Portugal María I. Arriortua Mineralogy and Petrology Department, Science and Technology Faculty, University of the Basque Country (UPV/EHU), Apdo, Bilbao, Spain Jose M. Barandiaran BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain; Department of Electricity & Electronics, University of the Basque Country, Bilbao, Spain Volodymyr A. Chernenko BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain; Department of Electricity & Electronics, University of the Basque Country, Bilbao, Spain; Institute of Innovative Research (IIR), Tokyo Institute of Technology, Yokohama, Japan; Ikerbasque, Basque Foundation for Science, Bilbao, Spain Daoyong Cong Beijing Advanced Innovation Center for Materials Genome Engineering, State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, China Guillermo Copello Universidad de Buenos Aires (UBA), Facultad de Farmacia y Bioquímica, Departamento de Química Analítica y Fisicoquímica, Junín, Buenos Aires, Argentina; CONICET
Universidad de Buenos Aires (UBA), Instituto de Química y Metabolismo del F
armaco (IQUIMEFA), Buenos Aires, Argentina V. Correia Centro/Departamento de Física, Universidade do Minho, Braga, Portugal; Centro ALGORITMI, Universidade do Minho, Campus de Gualtar, Guimar~aes, Portugal Carlos M Costa Centro de Física, Universidade do Minho, Braga, Portugal; Centro de Química, Universidade do Minho, Braga, Portugal Pedro Costa Centre of Physics of Minho and Porto Universities (CF-UM-UP), Campus de Gualtar, Braga, Portugal; IPC e Institute for Polymers and Composites, Universidade do Minho, Campus de Azurém, Guimar~aes, Portugal J.R. Dios GAIKER Technology Centre, Basque Research and Technology Alliance (BRTA), Parque Tecnol
ogico, Zamudio, Spain

xi

j

xii

Contributors

Roberto Fern
andez de Luis BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain Guillem Ferreres Departament d’Enginyeria Química, Universitat Politecnica de Catalunya, Terrassa, Spain Andreina García Department Water, Environment and Sustainability, Advanced Mining Technology Center (AMTC), Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Tupper, Santiago, Chile B
arbara Gonzalez-Navarrete Department Water, Environment and Sustainability, Advanced Mining Technology Center (AMTC), Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Tupper, Santiago, Chile Hideki Hosoda Institute of Innovative Research (IIR), Tokyo Institute of Technology, Yokohama, Japan Kristina Ivanova Departament d’Enginyeria Química, Universitat Politecnica de Catalunya, Terrassa, Spain Manfred Kohl Karlsruhe Institute of Technology, IMT, Karlsruhe, Germany R. Krishnapriya Sustainable Materials and Catalysis Research Laboratory, Department of Chemistry, Indian Institute of Technology Jodhpur, Jodhpur, Rajasthan, India Devika Laishram Sustainable Materials and Catalysis Research Laboratory, Department of Chemistry, Indian Institute of Technology Jodhpur, Jodhpur, Rajasthan, India Senentxu Lanceros-Mendez BCMaterials, Basque Center for Materials, Applications and Nanostructures, Parque Científico UPV/EHU Barrio Sarriena, Leioa, Spain; IKERBASQUE, Basque Foundation for Science, Bilbao, Spain Edurne Larrea (S.) Loire Valley Institute for Advanced Studies, Orléans & Tours, France; CEMHTI UPR3079 CNRS, Orléans, France Juan Manuel L
azaro-Martinez CONICET
Universidad de Buenos Aires (UBA), Instituto de Química y Metabolismo del F
armaco (IQUIMEFA), Buenos Aires, Argentina Kuntal Maity Organic Nano-Piezoelectric Device Laboratory, Department of Physics, Jadavpur University, Kolkata, West Bengal, India Sheila Maiz-Fern
andez BCMaterials, Basque Centre for Materials, Applications and Nanostructures, Leioa, Spain; Department of Physical Chemistry, Faculty of Science and Technology, University of Basque Country, Leioa, Spain

Contributors

xiii

Dipankar Mandal Institute of Nano Science and Technology, Mohali, Punjab, India P. Martins Centre/Department of Physics, University of Minho, Braga, Portugal; IB-S Institute of Science and Innovation for Sustainability, University of Minho, Braga, Portugal A. Gala Morena Departament d’Enginyeria Química, Universitat Politecnica de Catalunya, Terrassa, Spain J. Nunes-Pereira Centro de Física das Universidades do Minho e do Porto (CF-UM-UP), Campus de Gualtar, Braga, Portugal; Centre for Mechanical and Aerospace Science and Technologies (C-MAST-UBI), Universidade da Beira Interior, Covilh~a, Portugal Jo~ao Nunes-Pereira Centre of Physics of Minho and Porto Universities (CF-UM-UP), Campus de Gualtar, Braga, Portugal; Centre for Mechanical and Aerospace Science and Technologies (CMAST). Universidade da Beira Interior, Covilh~a, Portugal Joseba Orive Department of Chemical Engineering, Biotechnology and Materials, Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Beauchef, Santiago, Chile
Leyre Pérez-Alvarez BCMaterials, Basque Centre for Materials, Applications and Nanostructures, Leioa, Spain; Department of Physical Chemistry, Faculty of Science and Technology, University of Basque Country, Leioa, Spain Sílvia Pérez-Rafael Departament d’Enginyeria Química, Universitat Politécnica de Catalunya, Terrassa, Spain Yurieth Quintero (Marcela) Department Water, Environment and Sustainability, Advanced Mining Technology Center (AMTC), Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Tupper, Santiago, Chile Ander Reizabal-Para BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain B
arbara Rodriguez Department Water, Environment and Sustainability, Advanced Mining Technology Center (AMTC), Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Tupper, Santiago, Chile Maibelin Rosales Department of Chemical Engineering, Biotechnology and Materials, Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Beauchef, Santiago, Chile; Department Water, Environment and Sustainability, Advanced Mining Technology Center (AMTC), Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Tupper, Santiago, Chile

xiv

Contributors

Leire Ruiz-Rubio BCMaterials, Basque Centre for Materials, Applications and Nanostructures, Leioa, Spain; Department of Physical Chemistry, Faculty of Science and Technology, University of Basque Country, Leioa, Spain Paula G.-Sainz BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain; Mineralogy and Petrology Department, Science and Technology Faculty, University of the Basque Country (UPV/EHU), Apdo, Bilbao, Spain Daniel Salazar-Jaramillo BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain Jose Luis Sanchez Llamazares Instituto Potosino de Investigaci
on Científica y Tecnol
ogica A.C., San Luis Potosí, Mexico Rakesh K. Sharma Sustainable Materials and Catalysis Research Laboratory, Department of Chemistry, Indian Institute of Technology Jodhpur, Jodhpur, Rajasthan, India M. Shidhin The Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Chennai, Tamil Nadu, India M.M. Silva Department/Center of Chemistry, University of Minho, Campus de Gualtar, Braga, Portugal Gabriel Tovar Universidad de Buenos Aires (UBA), Facultad de Farmacia y Bioquímica, Departamento de Química Analítica y Fisicoquímica, Junín, Buenos Aires, Argentina; CONICET
Universidad de Buenos Aires (UBA), Instituto de Química y Metabolismo del F
armaco (IQUIMEFA), Buenos Aires, Argentina Carmen R. Tubio BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain Tzanko Tzanov Departament d’Enginyeria Química, Universitat Politécnica de Catalunya, Terrassa, Spain Kenji Uchino International Center for Actuators and Transducers, The Pennsylvania State University, University Park, PA, United States Ainara Valverde BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain José L. Vilas-Vilela (BCMaterials) Basque Centre for Materials, Applications and Nanostructures, Leioa, Spain; Department of Physical Chemistry, Faculty of Science and Technology, University of Basque Country, Leioa, Spain

Preface and Acknowledgments

All men by nature desire to know.

Aristotle (384 e 322 B.C.)

Multifunctional materials are increasingly being developed for technological applications including sensing and actuation, energy storage, production and conversion, water treatment, air purification, drug delivery systems, health monitoring, biomedical applications, structural health reinforcement, and monitoring, among others. Lightweight multifunctional materials are an emerging materials field for a next generation of multiresponsive and functional materials with improved performance and integration into devices, with an important role in the technological evolution to environmental friendlier and responsive materials contributing to interconnectivity and sustainability. In this book, the current state of the art on lightweight multifunctional materials of different nature, type, and morphologies is presented, as well as their preparation methods and applications. The book emphasizes the recent advances in these types of materials, including also their preparation and integration through novel processing paradigms such and additive manufacturing. In this scope, the book offers the first comprehensive account on this interesting and growing research field providing the main materials, effects and definitions, the current state of the art, the main research issues and challenges, and the actual and prospective application areas. All these are performed by selected authors with innovative and preponderant work in this field. The first chapter provides an overview to the area of lightweight multifunctional materials, the current state of the art of the different materials, types, and the manufacturing techniques and application areas, including multifunctional materials for energy generation and storage, sensing and actuation, self-healing, self-cleaning, antibacterial, thermochromic and electrochromic, and heating. Chapter 2 describes the additive manufacturing techniques used to produce multifunction materials, combining the intrinsic properties of the materials with electronic systems to integrate into functional devices. Chapter 3 describes porous lightweight materials for environmental remediation, with a focus on metaleorganic frameworks. xv

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Functional, lightweight materials: outlook, future trends, and challengesPreface and Acknowledgments

Their definition and requirements for metal ion recovery and their performance in the capture of heavy metals, among others, are addressed. Chapter 4 focuses on multifunctional materials for clean energy conversion, with special attention to nanomaterials with modified electronic, optical, or catalytic properties. Recent advances on materials for solar cells and fuel cells are presented in this chapter. Chapter 5 addresses the state of the art on thermochromic and thermoelectric materials, focusing on the main materials, effects, and applications. Chapter 6 presents multifunctional materials based on magnetic shape memory alloys with a focus on magnetic actuation and mechanical damping, while Chapter 7 describes piezoelectric polymers and composites as multifunctional responsive materials. Chapter 8 reports on the advances on electrochromic and electrorheological materials, considering also relevant aspects for applications such as layout of the devices and electrodes. Chapter 9 introduces high deformation multifunctional composites with a particular focus on structural health monitoring and self-sensing materials. Considering the increasing relevance of magnetic responsive materials, Chapters 10 and 11 present magnetoelectric composites and magnetoactive lightweight materials, such as magnetorheological, magnetostrictive, and magnetocaloric materials, respectively. Chapter 12 focuses on smart hydrogels for biomedical applications such as tissue engineering, soft robotic actuators, or controlled release of bioactive substances such as drugs, growth factors, or cells. Also, recent advances in the areas of 3D and 4D printing are presented. Chapter 13 focuses on antimicrobial lightweight materials with special attention to polymers, biopolymers, nanoparticles, hydrogels, foams, and surface coatings in order to increase antimicrobial properties. Finally, Chapter 14 summarizes some of the main open questions, challenges, and future trends of this research field. This book would not have been possible without the dedicated and insightful work of the authors of the different chapters. The editors really appreciate your acceptance to dedicate your precious time an effort to this adventure. We appreciate your kindness, dedication, and excellence in providing high-quality chapters that show the main characteristics, challenges, and applications of lightweight multifunctional materials. It was a pleasure and an honor to work with you on this important milestone in the area! This book could have not been also possible without the continuous dedication, support, and understanding from our research group colleagues

Preface and Acknowledgments

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both at the Center of Physics, University of Minho, Portugal, and the at BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain. Thank you all for the beautiful and continuous endeavor of driving science and technology a step further together and for sharing this important part of our lives! Last but not least, we truly thank the excellent support from the team from Elsevier: from the first contacts with Gwen Jones, Sumathi M. Sundaram, and Leticia M. Lima to the latter support from Surya Narayanan Jayachandran and Charlotte Rowley, passing through the different colleagues who supported this work. Your kindness, patience, continuous support, technical expertise, and insights were essential to make this book come true. It has been a real pleasure to work together with you! Finally, being the first book in this highly dynamic area, the author truly hopes that this work can become a milestone to further foster increasing scientific and engineering efforts and that lightweight multifunctional materials will increasingly become implemented as a new generation of high performance materials supporting a more sustainable and interconnected better world. Pedro Costa, Carlos M Costa, and Senentxu Lanceros-Mendez

CHAPTER ONE

Overview on lightweight, multifunctional materials Carlos M Costa1, 2 Pedro Costa1, 3, Senentxu Lanceros-Mendez4, 5 1

Centro de Física, Universidade do Minho, Braga, Portugal Centro de Química, Universidade do Minho, Braga, Portugal IPC e Institute for Polymers and Composites, Universidade do Minho, Campus de Azurém, Guimar~aes, Portugal 4 BCMaterials, Basque Center for Materials, Applications and Nanostructures, UPV/EHU Science Park, Leioa, Spain 5 Ikerbasque, Basque Foundation for Science, Bilbao, Spain 2 3

1.1 Introduction Multifunctional materials as smart materials that simultaneously maintain their structural functions increasingly over the years in both academic and industrial areas [1,2]. Novel developments in the overall properties of materials (physical and chemical modifications), processing methods, and integration into applications allow to specifically design materials for a wide variety of applications [2]. Considering the potential of these materials, the number of publications in this field has been strongly increasing in recent years, as shown in Fig. 1.1. Additive manufacturing and related manufacturing technologies allow, nowadays, the development and implementation of novel materials for a variety of applications, one area of interest being structural health monitoring (SHM), in which structural reinforcement in combined with nonstructural functions [2]. Multifunctional structural materials are materials that allow both structural function and non-structural function without the need for external devices. Structural function, including tailored strength, stiffness, fracture toughness, and damping, can be combined with nonstructural properties, such as the capability of providing noise and vibration control, electromagnetic interference (EMI) shielding, structural health monitoring, self-repair, thermo-chromism, self-cleaning, antibacterial, or energy harvesting/storage [2]. Multifunctional, lightweight materials are a concept allowing significant improvement in the overall efficiency of the total space system, improving performance and environmental friendliness [3]. Lightweight materials are Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00002-0

© 2021 Elsevier Ltd. All rights reserved.

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250 200 150 100 50 0 1970

1980

1990

2000

2010

2020

Years Figure 1.1 Number of publications about lightweight, multifunctional materials. Source: Scopus (keywords: Lightweight and polymer composites).

typically considered materials using polymers as host matrices or/and composites reinforced with fillers in order to address the specificity of each application or device. With polymer composites as material of choice, the intrinsic properties of the polymers (e.g., lightweight, flexibility, easy processing, etc.) can be combined with the unique properties of their reinforcements, such as electrical conductivity, high dielectric properties, or high mechanical stiffness and strength [4]. Polymer materials range from elastomers and thermoplastics to thermosets with a wide range of physicochemical properties that can be tailormade for applications. Further, in a broad sense, lightweight materials can also include inorganic high-density material that can be used in small amounts in lightweight structures to provide specific functions such as structural health monitoring or self-cleaning, without interfering in the lightweight characteristics of the overall structure. This chapter will provide an overview on the multifunctionality of polymer composites reinforced with specific fillers (from micro- to nanofillers with specific morphologies) for energy, shielding, sensor, or actuator materials. It does not pretend to be a comprehensive account of the field, which will be provided in the following chapters, but to highlight some of the main issues and concepts of this challenging and increasingly relevant field.

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1.2 Multifunctional materials classification and types Polymer-based materials capable of providing structural and nonstructural functions without the need of embedding or attaching external devices represent one of the main types of multifunctional structural material. This type of polymer-based materials can be classified by the following nonstructural functions (Fig. 1.2).

1.2.1 Energy generation and storage Energy generation or harvesting refers to the conversion of mechanical, thermal, or solar energy into electrical energy. Polymer-based materials for energy harvesting allow supplying energy to microdevices, replacing capacitors or batteries, reducing weight in the structures. Piezoelectricity [5], thermoelectricity [6], and dielectric elastomer generators [7,8] are some examples of electrical harvesting mechanisms to generate energy. Triboelectric effect can be described as the electrification of two different objects or materials during friction, resulting in a remarkable flow of electrons from one material to the other, balancing the potential difference between both materials. The most extensive polymers used in piezoelectric energy harvesting applications are the ones from the poly(vinylidene fluoride) (PVDF) family (mainly PVDF-TrFE and PVDF-HFP), due to their higher output performance [9] and physicochemical stability [9]. Conductive (carbonaceous or metallic) and ceramic nanofillers are also used as reinforcement materials, as well as blends with conductive polymers [9]. Carbon nanotubes and graphene as nanocarbonaceous particles, silver nanowires and zinc oxide particles as metallic particles, and barium titanate as ceramic are the most used materials for applications [9e11]. PVDF and copolymer or composites can generate some tens of microwatts of output

Energy (generation/storage) Sensors and actuators

Polymers +

Simple processing

Multifunctional materials

Self-healing, self-cleaning and antibacterial

Fillers Thermochromic and electrochronic Heating function

Others

Figure 1.2 Multifunctional, lightweight materials’ representative application areas.

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power for piezoelectric, 42 V output voltage for pyroelectric, and tens of microwatts in magnetoelectric devices [9]. The peak output power density of some triboelectric devices has been reported to be as high as 500 W/m2, through power conversion efficiencies up to 85% [12]. There are two main different ways to store electrical energy: capacitors or batteries [1]. Typically, capacitors are more used in high-power systems and batteries for lower-power ones. Among capacitors, supercapacitors are electrochemical devices without the need of anode or cathode, like batteries, but both of them need a separator to prevent the contact between the two electrodes and the electrolyte (ionic conductor). Regardless of the energy storage system, lightweight materials are used as separator/electrolyte placed between the electrodes, the most used types being single polymers, polymer composites, and polymer blends. There are different polymer types used for this application and the most widely used, from synthetic to natural polymers, are poly(propylene), PP poly(acrylonitrile) (PAN), poly(vinylidene fluoride) (PVDF), polyimide (PI), poly (ether-ether-ketone) (PEEK), cellulose, chitin, and poly(vinyl alcohol) (PVA) [9].

1.2.2 Sensing capabilities The sensing function is essential in lightweight materials for structural health monitoring, wearables, or materials for biomedicine [13]. As an example, in the area of structural health monitoring, strain/stress, bending, or damage is monitored in the structure by multifunctional materials with resistive, capacitive, inductive and magnetic, acoustic, and piezoelectric capabilities [4]. Resistive sensors are materials whose electrical resistance is affected by a particular physical quantity [4]. The resistance can vary as a function of changes in material properties, geometry, or a combination of both. Quantities typically measured using resistive effects include temperature, light, deformation, and magnetic field strength, among others. Tailoring the materials’ properties, resistive sensors have been also used for measuring force, torque, pressure, distance, angle, velocity, and acceleration [4]. In the following sections potentiometric sensors, strain gauges, piezoresistive and magnetoresistive sensors will be briefly discussed, as well as thermal and optical sensors based on resistive effects. Potentiometric sensors, strain gauges, piezoresistive sensors, magnetoresistive sensors, thermoresistive and light-dependent resistors (LDRs) are typical sensors to measure the distinct stimulus applied to a material.

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1.2.2.1 Potentiometric sensors Potentiometric displacement sensors can be linear or angular, consisting of a main material body, which is either wire wound or covered with a conductive film. In an ideal potentiometer, the voltage ratio (VR), defined as the ratio between the wiper voltage and the full voltage across the resistor is equal to x/L, where x is the distance between the initial and wiper positions and L is the electrical stroke. To measure the position of the wiper, the sensor is connected to a voltage source Vi with source resistance Rs; the output voltage on the wiper, Vo, is measured by an instrument with input resistance Ri. Ideally, the voltage transfer Vo/Vi equals the VR. Assuming that Rs ¼ 0 and Ri ¼ N the output voltage of the sensor is, respectively, for linear (Eq. 1.1) and angular (Eq. 1.2) potentiometer given by: x Vo ¼ Vi (1.1) L a Vi (1.2) Vo ¼ amax where L is the total electrical length and amax the maximum electrical angle. Potentiometer applications include linear and angular displacements and are widely used in all kinds of mechatronic systems to obtain information about angular positions of rotating parts. The potentiometers can also be used for the measurement of acceleration, force, pressure, and level [4]. 1.2.2.2 Piezoresistive sensors Piezoresistive sensors are based on the electrical resistivity variation of one material when this is deformed (tension, compression, or bending mode). This phenomenon was discovered in 1856 by William Thomson when observing the variation in resistance with elongation in iron and copper [14]. Numerous materials show piezoresistive effect, but only those with appropriate sensitivity are suitable to be applied in sensors [15]. Strain gauges [15], polymer-based composites [16], and semiconductor [17] materials are the materials used for sensing using the piezoresistive effect. Semiconductors germanium and silicon were the first materials widely used as piezoresistive materials, and reported the first measurements of large piezoresistive coefficients in these semiconductor crystals in 1954 [14]. Semiconductor materials show low mechanical properties (fragile, small deformations) with larger piezoresistive sensibility [14,18].

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The piezoresistive sensibility of the material can be determined by the Gauge Factor (GF), which can be expressed in relative resistance change per unit of strain: GF ¼

dR=R dl=l

(1.3)

where dR=R is the relative electrical resistance and dl=l is the relative deformation [16]. The GF has two main contributions in materials under strain [15]: GF ¼

dr=r þ 1 þ 2n dl=l

(1.4)

where the Poisson coefficient, n, of the material determines the geometric coefficient ð1 þ2nÞ and is always present in the sensors. The intrinsic resistance variation can be present in piezoresistive sensor or not, depending on the sample or type of measurement. An ideal rubber material shows a n ¼ 0.5, corresponding a GF ¼ 2, as the highest geometric factor in Eq. (1.4). Most metals or polymers show 0.2 < n < 0.35 [15,19]. Typically, the GF of strain gauges (metallic filaments deposited over a plastic substrate) changes from 2 to 2.5 and is larger (200) for semiconductor materials [4,15,19]. Piezoresistive sensors are suitable for the measurement of all kind of force-related quantities, for example, normal and shear force, pressure, torsion, bending, and stress. Polymer-based materials with large strain (elastomerlike matrices) are capable to measure from low to large strains, complementing the strain gauges’ typical range of strains, with large piezoresistive sensibility [20e23]. Piezoresistive sensors with strain larger than 50% and piezoresistive sensibility similar to semiconductor crystal have been achieved [24]. 1.2.2.3 Capacitive sensors A parallel plane capacitor is defined as an insulator material between two conductive electrodes. The dielectric properties of the insulator and geometry of both electrodes determine the capacitance (C) of the sensor, which for a parallel plate capacitor is given by: C ¼ ε0 εr

A d

(1.5)

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where ε0 permittivity of vacuum, εr is the dielectric constant of the material, A is area of the conductive electrodes, and d the thickness of the dielectric material. Capacitive sensors for displacement and force measurements have several advantages, being very robust and stable and applicable at high temperatures and in harsh environments. Capacitive displacement sensors commonly form part of instruments measuring pressure, sound, or acceleration [25]. The main advantages of lightweight materials in capacitive sensors are their flexibility and stretchability, easy preparation, and the possibility to be implemented in curved surfaces or unconventional geometries. Highly sensitive capacitive pressure sensors have been developed based on highly porous boron nitrate (BN)/polydimethylsiloxane composite foams (BNF@PDMS); this lightweight composite shows excellent mechanical resilience, extremely high compressibility (up to 95% strain), good cyclic performance, and superelasticity [26]. Carbon black/polydimethylsiloxane composites were produced for pressure sensors showing a sensitivity exceeding 35 kPa1 for pressures 10% for 40 wt.% PANI content), high electrical conductivity, 1 S/m, after the percolation threshold at z10 wt.% PANI and suitable piezoresistive response with gauge factor between z 1.5 and 2.4 for deformations up to 10% suitable for advanced electromechanical sensors applications. With the mixing of ceramic/polymers and nanosized active filler particles, the fabrication of polymer-derived silicate ceramics in various shapes and with variable composition, with potential applications in the fabrication

Additive manufacturing of multifunctional materials

Electrode reliability test specimen

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Electrical connection reliability test specimen

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(b)

(c)

(d)

Figure 2.3 Printed copper electrode: (a) photograph; (b) optical microscope top view; (c) cross-section view of the printed electrode; (d) experiment results of a scratch test using a diamond-tip pen. Reproduced with permission from H.S. Kim, J.S. Kang, J.S. Park, H.T. Hahn, H.C. Jung, J.W. Joung, Inkjet printed electronics for multifunctional composite structure, Compos. Sci. & Technol. 69 (2009) 1256e1264. https://doi.org/10.1016/j. compscitech.2009.02.034.

of components suitable for biological, high temperature, and functional applications, was achieved [25]. Three types of silicone resins (Silres MK, Silres H44 and Silres H62C, Wacker-Chemie GmbH, M€ unchen, Germany), two types of polysilazanes (Ceraset PSZ20, and perhydropolysilazane), and two types of fillers (TiO2 nanosized powder and Eu2O3 nanoparticles) were used. Shaping of the components was carried out using several plastic forming technologies, such as warm pressing, extrusion, injection molding, foaming, machining, fused deposition, and 3D printing. Another additive manufacturing technique, solvent-cast direct-writing, offers a low-cost, highly flexible and powerful fabrication route for microsystems featuring mechanical, microfluidic, and/or electrical functionalities, by using a thermoplastic polylactive polymer solution ink with dichloromethane [26]. Materials were processed by robotically controlled microextrusion of a filament combined with rapid solvent evaporation (Fig. 2.4).

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(a)

(b)

(c)

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2 mm

(e)

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Figure 2.4 Microstructures manufactured by solvent-cast direct-writing. (a) Top and side view virtual images of the programmed solvent-cast direct-writing fabrication of the square spiral. (b) Top and side view SEM images of a PLA square spiral. (c) Inclined top view SEM image of a PLA circular spiral. (d) Representative optical image of a PLA scaffold composed of nine layers. (e) Inclined top view SEM image of PLA 9-layer scaffold. (f) SEM image of a PLA cup. Reproduced with permission from S.Z. Guo, F. Gosselin, N. Guerin, A.M. Lanouette, M.C. Heuzey, D. Therriault, Solvent-cast three-dimensional printing of multifunctional microsystems, Small 9 (2013) 4118e4122. https://doi.org/10.1002/smll. 201300975.

Upon drying, the increased rigidity of the extruded filament enables the fabrication of complex freeform 3D shapes (high toughness fibers, spiral microchannels, and antennas). Melt electrowriting [27] with additive manufacturing principles was also developed. The experimental setup uses an electric field to generate a stable fluid jet with an expectable path that is uninterruptedly deposited onto a collector. The resulting fiber diameter is variable during the process and has potential for the biomedical field, providing a unique opportunity to perform low-cost, high-resolution, additive manufacturing research that is well positioned for clinical translation, using existing regulatory frameworks [28].

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Xu et al. [29] reported the development of 3D elastic silicone elastomerbased membranes shaped to match the epicardium of the heart via 3D printing, as a platform for deformable arrays of multifunctional sensors, electronic and optoelectronic components. Such devices completely envelop the heart, in a form-fitting way, and possess inherent elasticity, providing a mechanically stable biotic/abiotic interface during normal cardiac cycles. Semiconductor materials including silicon, gallium, arsenide, and gallium nitride co-integrated with metals, metal oxides, and polymers providing the multifunctional capability. Still in the field of membranes, Gao et al. [30], presented a combined inkjet printing and template synthesis technique to prepare charge mosaic membranes in a rapid and straightforward manner, demonstrating the unique transport properties that result from the mosaic membrane design. For that, poly(vinyl alcohol)-based composite inks containing poly(diallyldimethylammonium chloride) or poly(sodium 4styrenesulfonate) were used to pattern positively charged or negatively charged domains, respectively, on the surface of a polycarbonate tracketched membrane with 30 nm pores. The developed membranes can be deployed in nanoscale technologies that rely on the selective transport and separation of ionic solutes from solution. In a theoretical study [31] by numerical simulations it was discussed that the multifunctional response of polymeric systems (acrylonitrile butadiene styrene and polylactic acid) can be manipulated by axial grading and that optimal design/fabrication of multifunctional smart structures by 3D printing may be developed for vibration control applications. Most importantly, such modeling and resulting solutions framework can also be used for analyzing the vibration of the structures, showing that smart materials can be incorporated for axial grading and may be manipulated for vibration control applications. Polylactic acid was also used by Prashantha et al. [32] as a matrix for graphene aiming to allow additive multilayer deposition of the polymeric multifunctional nanocomposite. The resulting 3D printed polylactic acid/ graphene nanocomposites containing 10 wt.% graphene in PLA matrix exhibited improved mechanical and thermo-mechanical properties, becoming a promising new 3D printable biomaterial for tissue engineering, bioelectronics, and biosensors. Keeping the focus on polymer-based graphene nanocomposites, Yamamoto et al. [33] addressed the problems of fused filament fabrication or fused deposition modeling, namely the low strength and toughness of the typical thermoplastic materials such as acrylonitrile butadiene styrene, polylactic

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acid, and polybutylene terephthalate. For that, the team introduced 0.06 wt.% of graphene oxide on an acrylonitrile butadiene styrene matrix that led to a much higher strain-to-failure (14% for facedown and 29% for upright) and toughness (20% for facedown and 55% for upright) while also increasing the fracture strength (3.5% for facedown and 10% for upright) and decreasing the stiffness (6% for facedown and 15% for upright). Such improvements in strain-to-failure, toughness, and fracture strength have shown the multifunctionality of the developed printed system where several mechanical properties are improved for a number of structural applications. Unique properties including controlled surface slipperiness, self-reporting on the loss of liquid repellency, and sensing the temperature of contacting liquids demonstrated on the printed nanocomposites with lubrication treatment were reported on ultrahigh-molecular-weight polyethylene/SiO2 nanocomposites Fig. 2.5 [34]. The functional printing of ultrahigh-molecular-weight polyethylene/ SiO2 nanocomposites was performed by using a commercially available inkjet printer (L801, Epson, Japan) and shows applicability in antifouling coatings, food/medical packaging, smart windows, and sensors. While many “ad hoc” designs of 4D printed solutions have been progressively developed for specific processes, the general approach to produce smart materials by additive manufacturing techniques, in real time across an entire product development process, is not pervasive in the industry. To solve this issue, Wenqing et al. [34] proposed a general 4D printingoriented framework for the design of multifunctional, shape-memory polymer architectures. Such report was not intended to be an exhaustive and specific instruction but is instead a means to motivate to seek the process for applying these unique functional materials for specific designs and applications. Multiwall carbon nanotubes were found to improve the electrical and dielectric properties, to promote ultrahigh polarization density, and to form local microcapacitors within poly(vinylidene) fluoride/BaTiO3 composites [35]. Additionally, the 3D printing process of those materials provided homogeneous dispersion of nanoparticles, alleviating agglomeration of nanoparticles, and reducing the microcracks/voids in the matrix. Such promising results opened the way for the 3D printing of multifunctional nanocomposites with temperature and strain sensing capabilities, increased mechanical property, as well as the feasibility for large-scale multifunctional sensor device manufacturing with freedom of design and low cost.

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Figure 2.5 (a) Schematic representation of print-assisted functionalization on porous nanocomposites for multifunctional liquid-infused materials. (b) Surface morphology of UHMWPE/SiO2 nanocomposites with different SiO2 contents. (c) Optical images of the corresponding UHMWPE/SiO2 nanocomposites after inkjet printing. (d) Crosssection of Printed-S0, Printed-S3.5, and Printed-S7 samples showing distinct ink penetration depth. Scale bar: (b) 5 mm; (c) 300 mm; (d) 200 mm. Reproduced with permission from W. He, P. Liu, J. Jiang, M. Liu, H. Li, J. Zhang, Y. Luo, H.Y.Cheung, X. Yao, Development of multifunctional liquid-infused materials by printing assisted functionalization on porous nanocomposites, J. Mater. Chem. 6 (2018) 4199e4208. https://doi.org/10.1039/ c7ta10780c.

By mixing materials that are simultaneously electroactive and magnetoactive, and knowing that the successful application of those magnetoelectric (ME) materials is closely related to the processing and integration of additive manufacturing techniques, Lima et al. [36] developed novel screen-printed and flexible ME materials composed of poly(vinylidene fluoride-co-trifluoroethylene) P(VDF-TrFE) as the piezoelectric phase and poly(vinylidene fluoride) (PVDF-CFO) as the magnetostrictive one. Such all-printed ME composite exhibited an ME voltage coefficient (a) of 164 mV cmOe1 at a longitudinal resonance frequency of 16.2 kHz. The optimized magnetic, piezoelectric, and ME behavior, together with the reduced cost of assembly, easy integration into devices, and the

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possibility of being obtained over flexible and large areas through additive manufacturing techniques demonstrated the suitability of the advanced lightweight, multifunctional materials for applications in printed electronics, sensors, actuators, and energy harvesters (Fig. 2.6).

2.2.2 Materials for electronics Electronic systems, and the whole area of electronics in general, have shown exponential growth in recent decades, as demonstrated by a simple technological comparison between the current date and the past decade. This evolution is supported by a correlated growth in advanced materials, with particular focus in the last decade on (multi)functional materials. This combination of extremely efficient and miniaturized electronic systems and everperforming functional materials makes it possible to move to new and exciting research scenarios. While the interaction of these two worlds is not yet easy, the evolution of additive manufacturing systems allows this combination to become increasingly easy, thus creating solutions previously imaginable in terms of integration, flexibility, size, and autonomy. By examining the domains of additive manufacturing within the electronics industry, it is realized that inkjet printing techniques are increasingly being used to produce electronic systems, including circuit boards. Typically, this method involves a printhead that works on a horizontal surface by applying conductive inks that allow the rapid production of custom circuit boards. However, a big revolution is expected and corresponds to the complete manufacturing of electronic circuitry, electronic components and the structural part of the device, resulting in a fully functional product digitally fabricated by purely additive processes. Thus, research in this field is further developing this concept, where specific works are pointing out this way, as it is the case of printed organic transistors (OTFTs), reported since 2002 [37]. Most of these reports make use of flexible substrates, but the fabrication is based on subtractive technology. The fabrication of fully printed OTFTs has been also reported, such as in Castro et al. [38], where a bottom-gate approach is used to print fully functional four-layer inkjet-printed OTFTs, the device being characterized by an electron mobility of 0.012 cm2/Vs1 and on/off ratio of 103. More recently, higher performance has been achieved by improving both active materials and printing technologies, the actual main challenge being the manufacturing success rate versus device efficiency [39], as only then the devices will become interesting for industrial applications.

thermoplastic, for 3D multifuncional architecture [32][35]. multifuntional material whith magnetoelectric response for magenetica and magnetoelectric sensores [36]. Polymer binder for high-performance printable lithium-lon battery electrodes[23], and sensor [24]. multifunctional shape memory polymer based for multifuncional architecture [34]. Viscoelastic materials for tissure enginneering [30][31]. Poly(vinyl alcohol)-based ink for charge mosaic membranes- 30 nm pores resolution [29]. 3D elastic silicone elastomer-based membranes for deformable arrays of multifuctional sensors[28]. Thermoplastic polymer-based for 3D microsystens - resolution 45Pm[26]. Active fillers into preceramic polymers for multifunctional ceramic components, resolution 0.4mm [25].

320 µm 160 µm 80 µm

Copper nano-ink based multifunctional comosite structures for electronic aplications, resolution 80Pm [22]. Multifunctional bio-polymer for cell microarry aplication [21].

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Figure 2.6 Schematic summary of the evolution of the development of multifunctional materials through additive manufacturing techniques. 35

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Another major field of development in the field of printed electronics involves the printing of OLEDs where one of the solutions developed to maximize performance is the use of light emitters, inorganic or hybrid materials, such as inorganic quantum dots (QDs), and inorganic fluorescent dyes. Singh et al. [40] demonstrated the use of a hybrid organic-inorganic material and inkjet printing for the fabrication of the emitting layer. It was shown that the device exceeded 10 kcd/m2 for rigid substrates and 9.6 kcd/m2 for flexible substrates, respectively. A major reason for using LEDs is the increasing demand of digital displays. In this field, Haverinen et al. [41] presented displays manufactured using QDs of different sizes on a full-control red, green, and blue direct current graphics matrix with a brightness of 100 cd/m2. The only disadvantage of using inorganic QD materials is their high cost [42]. Thus, carbon nanotubes (CNTs) are being explored also for developing display devices. Shigematsu et al. [43] demonstrated printed electrodes based on single-wall carbon nanotubes coated with phosphor as the counter-electrode for emissive displays fabricated using electrostatic inkjet technique. The study and development of passive electronic components have also proven to be areas of strong research and need, always with the ultimate goal of achieving fully printed devices. V. Correia et al. [44] proved that it was possible to design and print resistances, capacitors, and coils with specific characteristics, solely through inkjet printing techniques [45]. The feasibility of stacking these devices has also been explored and demonstrated to increase the efficiency of the end devices per unit area, thus competing with traditional manufacturing methods. Research on printed sensors is another major research focus, with physical sensors such as pressure, strain, magnetic field, and current, among others already being reported [46]. Nevertheless, efforts have to be devoted to meet industrial demands, mainly due to the lack of functional materials capable of guaranteeing reproducibility and durability over time, a factor which the new generations of functional materials already is addressing [47]. Finally, another high impact area, where once again functional materials are among the most suitable solution, corresponds to the energy area, with the batteries being printed along with the capacitors, the devices that by far show the best results. Gaikwad et al. [48] demonstrated high-potential fully printed batteries with polyvinyl alcohol and cellulose being used as a substrate and separator and KOH and ZnO solution as electrolyte. The electrodes were printed by stencil printing and based on zinc and MnO2. As a

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result, the dimension of the printed pattern depends on that of the stencil’s mesh. The results showed that the initial open circuit potential of the entire battery was 14 V. Currently some companies in the traditional battery business already have fully printed market solutions [49]. From the presented overview the strong development of printed electronics based on lightweight materials and the real dependence on the evolution of functional materials, as well as their implementation in industrial manufacturing processes are confirmed.

2.3 Future trends It has been shown that advanced lightweight, multifunctional materials have already found uses in a large variety of technological devices, being the inherent flexibility of additive manufacturing technologies to fabricate complex geometries with spatially varying distributions of phases that can be engineered to tailor mechanical and physical properties in a precise way [50]. Nevertheless, and in a general way, such a concept is still in its first steps. As it matures, and in an exciting Internet of Things context, it will be part of daily life. While smart materials science has typically focused on the development of functional materials based on inorganic components [19], it is important to take into account that there are an increasing number of lightweight materials that have also shown multifunctional behavior. Nevertheless, and despite the advantages of combining the Internet of Things, the multifunctional smart materials, and printing processing concepts, most of the current commercially available devices are not based on printing technologies. This multidisciplinary concept has a strong scientific and economic potential in this increasingly interesting field. Considering their well-known advantages such as fast curing processes at room temperature with reduced volatile organic compound emissions, UV curable multifunctional materials are becoming of general research and implementation interest. Chromic, self-healing, shape memory, piezoresistive, or piezoelectric materials are just a few examples of them. Parameters such as viscosity, density, surface tension, and contact angle on different substrates are needed to be discussed/optimized in detail when these multifunctional materials are reported, once they strongly influence their processability and integration into applications. Furthermore, depending on the nanofillers employed, the multifunctionality can be also affected. Thus, obtaining optimized materials in terms of functionality/smart response as well as processability requires a deep understanding on the filler-polymer

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interaction. In this way, the potential solutions and/or new approaches that are being investigated include the use of polymer-coated fillers, new filler dispersion techniques, or new fillers more compatible with UV-curable resin (i.e., that are not affected by and do not affect the UV curing process) [51]. It is also important to notice that additive manufacturing technologies can be used to produce a large combination of alloys, metals, ceramics, and composites in different geometries. It is, however, this same flexibility that renders additive manufacturing technologies difficult to develop into reliable commercial products. Additionally the additive manufacturing optimization generally does not carry out from one process to another, making it very difficult to generalize operational principles for the various additive manufacturing technologies for the production of multifunctional materials, such as would be required by industry for proper operation of a reliable manufacturing line [50]. In the next years, 4D printing techniques will allow an innovative and disruptive effect in this field due to the quality, efficiency and performance of this technique. In the particular case of biomedical sciences, 4D printing will allow customized production for each individual patient, whose smart implants, tools, devices, organ printing, tissue engineering, and selfassembling human-scale biomaterials can be easily achieved in less time, which has extensive benefits to the patients [50]. Such approach will replace the conventional and limited scaffold production methods, leading to new possibilities in the biomedical field. Additive manufacturing on the production of photodetectors and UV curable polymer-based multifunctional materials will allow applications for consumer electronics, with added advantages in their production such as fast curing at room temperature, space and energy efficiency, highresolution patterns, and solvent-free formulations [51]. The most interesting is that many of the applications or products related with multifunctional materials, where additive manufacturing can have a huge impact, are not listed in this chapter as they do not currently exist, as traditional engineering has limited adaptability/ability to meet the current design demands. As the hardware, software, and materials capabilities in additive manufacturing tailored for the production of multifunctional materials continue to develop, new materials, new architectures, new geometries, and smart 3D/4D multifunctional objects will present new, challenging, and exciting opportunities in the near future. Everything else is a . challenging innovation roadmap.

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Acknowledgments The authors thank the FCTdFundaç~ao para a Ciência e Tecnologiadfor financial support in the framework of the Strategic Funding UID/FIS/04650/2020 and under project PTDC/ BTM-MAT/28237/2017 and PTDC/EMD-EMD/28159/2017. P. Martins (CEECIND/ 03975/2017- assistant researcher contract Individual Support e 2017 Call) and V. Correia (DL57/2016 junior researcher contract) thank FCT for the contract under the Stimulus of Scientific Employment. The authors acknowledge funding from the Basque Government Industry and Education Department under the ELKARTEK, HAZITEK, and PIBA (PIBA-2018-06) programs, respectively. Funding from the European Union’s Horizon 2020 Program for Research, ICT-02-2018 - Grant agreement no. 824339 e WEARPLEX is also acknowledged. The authors thank funding by the Spanish State Research Agency (AEI) and the European Regional Development Fund (ERFD) through the project PID2019-106099RB-C43/AEI/ 10.13039/501100011033.

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[40] M. Singh, H.S. Chae, J.D. Froehlich, T. Kondou, S. Li, A. Mochizuki, G.E. Jabbour, Electroluminescence from printed stellate polyhedral oligomeric silsesquioxanes, Soft Matter 5 (2009) 3002e3005, https://doi.org/10.1039/b903531a. [41] H.M. Haverinen, R.A. Myllyl€a, G.E. Jabbour, G.E. Jabbour, Inkjet printed rgb quantum dot-hybrid led, IEEE/OSA J. Disp. Technol. 6 (2010) 87e89, https://doi.org/ 10.1109/JDT.2009.2039019. [42] N. Saengchairat, T. Tran, C.K. Chua, A review: additive manufacturing for active electronic components, Virtual & Phys. Prototyp. 12 (2017) 31e46, https:// doi.org/10.1080/17452759.2016.1253181. [43] S. Shigematsu, Y. Ishida, N. Nakashima, T. Asano, Electrostatic inkjet printing of carbon nanotube for cold cathode application, Jpn. J. Appl. Phys. 47 (2008) 5110e5112, https://doi.org/10.1143/JJAP.47.5109. [44] V. Correia, K.Y. Mitra, H. Castro, J.G. Rocha, E. Sowade, R.R. Baumann, S. Lanceros-Mendez, Design and fabrication of multilayer inkjet-printed passive components for printed electronics circuit development, J. Manuf. Process. 31 (2018) 364e371, https://doi.org/10.1016/j.jmapro.2017.11.016. [45] H.F. Castro, V. Correia, E. Sowade, K.Y. Mitra, J.G. Rocha, R.R. Baumann, S. Lanceros-Méndez, All-inkjet-printed low-pass filters with adjustable cutoff frequency consisting of resistors, inductors and transistors for sensor applications, Org. Electron. 38 (2016) 205e212, https://doi.org/10.1016/j.orgel.2016.08.025. [46] D.E. Schwartz, Flexible electronics systems and advances, Proc. SPIE Int. Soc. Opt. Eng. (2019). [47] S.F. Anis, R. Hashaikeh, N. Hilal, Functional materials in desalination: a review, Desalination 468 (2019), https://doi.org/10.1016/j.desal.2019.114077. [48] A.M. Gaikwad, A.C. Arias, D.A. Steingart, Recent progress on printed flexible batteries: mechanical challenges, printing technologies, and future prospects, Energy Technol. 3 (2014) 305e328, https://doi.org/10.1002/ente.201402182. [49] Report ID: GVR-3-68038-526-7, Flexible, Printed and Thin Film Battery Market Size, Share & Trends Analysis Report by Battery Type (Rechargeable, Disposable), by Application (Consumer Electronics, Smart Cards, Medical), by Region, and Segment Forecasts, 2019e2025, 2019. [50] C. Zhang, F. Chen, Z. Huang, M. Jia, G. Chen, Y. Ye, Y. Lin, W. Liu, B. Chen, Q. Shen, L. Zhang, E.J. Lavernia, Additive manufacturing of functionally graded materials: a review, Mater. Sci. & Eng. A 764 (2019) 138209, https://doi.org/10.1016/ j.msea.2019.138209. [51] C. Mendes-Felipe, J. Oliveira, I. Etxebarria, J.L. Vilas-Vilela, S. Lanceros-Mendez, State-of-the-art and future challenges of Uv curable polymer-based smart materials for printing technologies, Adv. Mater. Technol. 4 (2019) 1800618, https://doi.org/ 10.1002/admt.201800618.

CHAPTER THREE

Porous, lightweight, metal organic materials: environment sustainability Ainara Valverde1, Paula G.-Sainz1, 8, Joseba Orive2, Edurne Larrea (S.)3, 4, Ander Reizabal-Para1, Gabriel Tovar5, 6, azaro-Martinez6, Guillermo Copello5, 6, Juan Manuel L
7 B
arbara Rodriguez , B
arbara Gonzalez-Navarrete7, Yurieth Quintero (Marcela)7, Maibelin Rosales2, 7, Andreina García7, María I. Arriortua8, Roberto Fern
andez de Luis1 1

BCMaterials (Basque Centre for Materials, Applications & Nanostructures), Bld. Martina Casiano, Leioa, Spain Department of Chemical Engineering, Biotechnology and Materials, Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Beauchef, Santiago, Chile 3 Le Stusium Research Fellow, Loire Valley Institute for Advanced Studies, Orléans & Tours, France 4 CEMHTI - UPR3079 CNRS, 1 Avenue de la Recherche Scientifique, Orléans, France, France 5 Universidad de Buenos Aires (UBA), Facultad de Farmacia y Bioquímica, Departamento de Química Analítica y Fisicoquímica, Junín, Buenos Aires, Argentina 6 CONICET
Universidad de Buenos Aires (UBA), Instituto de Química y Metabolismo del F
armaco (IQUIMEFA), Buenos Aires, Argentina 7 Department Water, Environment and Sustainability, Advanced Mining Technology Center (AMTC), Facultad de Ciencias Físicas y Matem
aticas, Universidad de Chile, Av. Tupper, Santiago, Chile 8 Mineralogy and Petrology Department, Science and Technology Faculty, University of the Basque Country (UPV/EHU), Apdo, Bilbao, Spain 2

3.1 Introduction Metal organic frameworks (MOFs) are crystalline inorganic-organic hybrid materials that are assembled and sustained by coordination bonds between metal ions or metal-oxo clusters and negatively charged organic linkers bearing a complexing function (e.g., carboxylate, phosphonate, azolate .) [1,2]. The emerging research interest in MOFs results from the simultaneous occurrence of five key characteristics that are (i) the crystallinity, (ii) the tunable porosity, (iii) high surface area (MOF materials have so far the highest surface area of all porous materials), (iv) the existence of strong metal-ligand interactions, and (v) the structural diversity that allows to precisely design materials for a particular application (Fig. 3.1). All these

Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00012-3

© 2021 Elsevier Ltd. All rights reserved.

43

j

44

Synthesis

Organic ‘link’ Metal oxide ‘joint’

(a)

(b)

(c) Ainara Valverde et al.

Figure 3.1 Schematic representation of Metal Organic Framework components (a), their assembly in a porous ordered structure (b) and their functionalization at the inorganic and organic subunits.

Porous, lightweight, metal organic materials: environment sustainability

45

aspects together with the pores chemical functionalization have made MOFs ideal candidates for adsorption and separation applications both in gas and liquid matrixes. Specifically, MOFs have shown an outstanding performance over metal ion recovery from different aqueous media. Once established the main chemical MOF modification paths to tune the hosteguest chemistry over target gases and organic molecules; it was a natural step for the scientific community working on MOFs to apply the previously acquired knowledge to metal ion capture. In the first attempts, MOFs were mainly used for water depollution purposes in model laboratory solutions. It was promtly demostrarted the potentials of MOFs for this end, overcoming the performances of classic adsorbents over the capture of cationic lead, mercury, cadmium, or anionic arsenate, selenite, or chromate species. Some of these works applied multi-experimental and computational approaches to unravel the underpinning mechanisms governing the metal ion captures in MOF materials; being the local structure of the adsorbed species and metal ion docking points within the MOF networks determined in different extents of accuracy. Once confirmed the feasibility of MOFs, and concretely the chemical lability and plasticity of their inorganic and organic components to be functionalized with increasingly complex metal trap-chelating motifs [3,4], MOFs have been applied successfully in more stringent and complex media, in terms of acidity, ionic strength, multicomponent metal ions mixture, or even radioactivity. Depending on the application scenario, MOF chemistry has shown its feasibility to recover and separate metal ions of different natures and charges from acid waters, multi-element metal leachates, seawater, radioactive aqueous waste or electronic scratch acid solutions, just to mention some of the most relevant scenarios. In addition to high capacities, some MOFs have demonstrated: (i) tailorable pore surface chemistry that can be directed toward the specific adsorption of low concentrated ions in the presence of high concentrated competitor species, (ii) to work in a repeatable manner; (iii) enough chemical stability, at least to be used in proof-of-concept studies, without a dramatic loss of their performance over cycling. In this chapter, an overview of the chemical and structural functionalization paths of MOFs applied to metal ions recovery has been summarized, taking special care of the keyelock interactions derived from the chemical characteristics of the metal chelating agents grafted within MOFs, and the chemical nature of the adsorbed target metal ions. Despite most of the developed studies have applied an strategy based on: “one MOF functionalization

46

Ainara Valverde et al.

applied to the adsorption of one metal ion” it is worthy to mention that merging all the information of the different researches, a general landscape of the fundamentals of the MOF chemistry for metal ions recovery can been derived.

3.2 Requirements of MOF materials for metal ion recovery Applicability of MOFs for metal ion recovery/separation at different scenarios requires from several combined characteristics that the adsorbent need to accomplish, the most important being: (1) Long-term chemical stability during operation (2) High adsorption capacity (3) Fast kinetics (4) High selectivity over target metals, even in the presence of competitor ions (5) Stability to the regeneration during the elution of the adsorbed species and negligible capacity fading along the reusability cycles of the adsorbent Depending on the application scenario and its final objective (Table 3.1), the weight of each of the above mentioned adsorbent characteristics differs. For example, heavy metal decontamination of polluted water does not requires a high selectivity toward single metal ions, but from an efficient removal of target heavy metals at low concentrations independently, if side adsorption of other metal species occurs in parallel over the process. On the contrary, the adsorbents that are applied to the recovery of highly valuable and technologically relevant metals form acid leachates (e.g., radioactive waste, electronic scratch recycling, REE separation, etc.) need to meet excellent chemical stability, high adsorption capacity, and outstanding specificity toward single ions in the presence of highly concentrated competitor ions. Therefore, the higher the selectivity over the target ions, the less the number of adsorption/desorption runs are needed until the desired purity range of the target element is achieved [1,2]. Meeting these requirements on the same material is not straightforward; and even less if it is considered that the powdered nature of MOFs makes it necessary to shape them as mechanically stable macroscopic devices (membranes, filter, etc.) to use them close to real-application conditions [3]. (6) Regeneration and reusability of the adsorbent

Porous, lightweight, metal organic materials: environment sustainability

47

Table 3.1 Potential application fields of metal organic framework materials for metal ion recovery from different scenarios. Applicability fields Chemical characteristic

Polluted water sources

Sea water

Industrial and mining acid drainage waters

Acid leachates

Radioactive waste waters Bioleaching

Fossil fuels or organic solvent purification

• Acidity: pH 7epH 5 • Ionic strength: low to moderate. • Composition of water sources strongly depends on the geological environment to which are exposed the surface and underground water streams. • Acidity: pH 7epH 8 • Ionic strength: high • Presence of high concentrations of sodium, chloride, sulfate ions. • Acidity: pH below 3 • Ionic strength: moderate to high • Presence of highly chelating inorganic species, such as sulfate and phosphate groups. Composition strongly depends on the chemical nature of mine tailings. Possible presence of some radioactivity level (e.g. phosphor-gypsum deposits). • Acidity: pH between 0 and 3. • Ionic strength: very high. • High to low concentration of the target metal ions. This depends on the composition of end-of-life products dissolved in concentrated acids (e.g. permanent magnets, phosphorous lamps, electronic scratch, batteries). • Acidity: pH between 1 and 3. • Ionic strength: very high • Radioactive media. • Acidity: pH between 1 and 4. • Ionic strength: very high • Presence of organic molecules that could disrupt the metaleligand framework connectivity leading to the MOF degradation. • Hydrophobic media. • Presence of organic molecules that could disrupt the metaleligand framework connectivity leading to the MOF degradation.

48

Ainara Valverde et al.

We have to mention that MOF shaping is not the main scope of the chapter, a brief overview has been given at the end of the chapter in order to envisage the future trends in this direction, since an effective supporting of active adsorbents that keep active its full capacity and specificity is a critical issue to be addressed to apply them in close to real-application conditions [4].

3.2.1 Chemical stability assessment Water and chemical stability of MOFs mainly depend on the degree of metal-ligand bonds hydrolysis during the operation conditions [5]. Specifically, the strength of the metal-ligand bond is directly related to the acidity of the metal at the inorganic clusters, and the basicity of the chelating groups at the organic linkers. As a general rule, the higher the charge density (Z/r2) of the metal ions at the inorganic clusters (e.g., small radii high valence elements such as Zr(IV), Ti(IV), Cr(III), Al(III)[120]), the better the chemical resistance to water of the metal-organic frameworks [6]. Indeed, following this approach, some of the benchmark Zr(IV), Ti(IV), Cr(III), and Al(III) [7] water, and even acid and basic resistance MOF materials have been crystallized during the last decade [8]. In addition to the formation of strong metaleligand bridges, other factors such as (i) nuclearity of the inorganic units, their (ii) connectivity degree through the organic linkers, and the (iii) chemical nature of the organic linkers itself, influx drastically the chemical stability window of MOF materials. Within the MOFs composed from the same metal ions, those exhibiting inorganic building units with higher nuclearities, such as Fe-MIL53 [9], Zr-MIL140 [10], or Ti-MIP177-HT [11], which contains linear inorganic oxo-chains within their crystal structures, show higher chemical resistance to acid and basic media than MOF compounds formed from isolated inorganic building units containing the same metal (Fe-MIL88 [12], Zr-UiO66 [13], or Ti-MIL125 [14]). The chemistry of the organic linkers, as well as their geometry, and connection points to the inorganic building units also affect the chemical resistance of the MOFs. Therefore, organic bridges with higher anchoring points, tetrapodal > tripodal > bipodal, usually generate MOFs with a wider chemical applicability window in terms of pH value. In addition, the same inorganic and organic building blocks can be arranged in different MOF polymorphic structures with varied topologies, and hence connectivity; which exhibit specific chemical resistance in function of the connectivity and chemical robustness of the components.

Porous, lightweight, metal organic materials: environment sustainability

49

The chemistry of the organic ligand also influx the chemical stability of the extended frameworks, since MOFs formed from strong and rigid moieties, such as benzene bearing linkers exhibit higher chemical resistance than aliphatic chainebased ones, even if the connectivity and topology of the MOFs’ net is the same. One of the most clarifying examples in that regard is the varied chemical resistance shown by the derivatives of UiO-66 framework obtained with different decoration motifs at the organic linkers (e.g., -NH3, -NO2, -(COOH)2, -(SH)2), or with organic building blocks of increasingly lengths (i.e., UiO66 < UiO67 < UiO68). Taking into consideration all the aspects concerning the chemical and water stability of MOF frameworks, some general trends have been summarized in Fig. 3.2. For more detailed information the readers are referred to the book chapter recently published by Mouchaham et al. [15]. As a general rule, MOFs containing high valence metal ions, such as Ti(IV), Zr(IV)-Ce(IV)-Hf(IV), Cr(III), Al(III), REE(III), and some Fe(III) carboxylate MOFs exhibit enough chemical stability to perform initial proof-of-concept studies in aqueous solutions. The applicability of REE(III) and Fe(III) MOFs is restrained in terms of acidity or basicity of the media in comparison to Zr(IV), Cr(III), Al(III) or Ti(IV), the pH stability window of which covers a wider application range Fig. 3.2. Despite some divalent-carboxyl, azolate, and imidazole based MOFs are quite effective for water decontamination, ’their operational window has to be assessed in terms of the stability on different chemical media individually for each case. As an archetype example, ZIF-8 material, formed from Zn(II) and methyl-imidazole linkers, shows an outstanding stability at neutral or basic conditions, but it dissolves partially or completely at acidic media.

Chemical stability in acid media Cr(III)-carboxylates M(IV)-phosphonates Stable in concentrated acid solutions (pH0 – 1– Cd(II) > Hg(II) sequence. As softintermediate acids, their chelating molecules counterpart are usually formed by soft to intermediate basic motifs exhibiting intermediate to low pKa values. Following this general rule, MOFs for mercury, lead, and cadmium capture have been decorated with amine, amide, thymine, thiol, and other

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms. Target metal/capacity (mg/g)/equilibrium Compound Application conditions Adsorption mechanisms time

Reusability

Refs.

Cd(II), Pb(II), Hg(II) NR

Not reported

[45]

Hg(II)/950 mg/g/20 min 250 mg/g/20 min

Selectivity studies over [46] other ions usually present in water, such as Naþ, Kþ, Ca2þ, Mg2þ, HCO3, Cl, and NO3 High selectivity and [47] ultrafast kinetics Not reported [48]

CATIONS Weak and intermediate acidic heavy ions (Hg, Cd, Pb)

AMOF

Single element solution

CaCu-Methionine BioMOF

Single element solution

Cu-Methionine BioMOF Cr-MIL-101Alkenylthiol DUT-67

Single element solution

Cation exchange with organic cations located at the pores of the structure Coordination to methionine fragments within the MOF pores. Speciation determined by single crystal X-ray diffraction after adsorption

Single elements solution

Thiol e Hg(II) interaction

Hg(II)/NR/NR

2,5-thiophenedicarboxylate eHg(II) interaction

Hg(II)/NR/30 min

FJI-H12

Single element solutions and mercury polluted real water samples Single-element solutions

S¼C]N moieties as active Hg(II) chelators

Hg(II)/439 mg/g/2 h

HKUST-Dithioglycol

Single-element solutions

Dithioglycol

Hg(II)/714 mg/g/200 min

Effective mercury [37] removal from real water samples Regeneration with a [49] KSCN solution induce a capacity loss. Regeneration not [24] tested (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium Compound Application conditions Adsorption mechanisms time

Reusability

Refs.

UiO-66-SH2

Not reported

[50,51]

Single-element solutions

Cr-MIL-101 e Thymine Single-element and real water samples

MOF-5-SH2

Single-element solutions

FJI-H9

Single-element solution

Proved SH$Hg(II) affinity Hg(II)/225 mg/g/NR derived from adsorption capacity results. Direct S-Hg coordination derived from IR spectroscopy Hg(II)/50 mg/g/ Thymine can coordinate 20 min with Hg2þ effectively

Dimethyl 2,5dimercaptoterephthalate Dimethyl ammonium cation exchange

Hg(II)/NR Hg(II)/275 mg/g/NR Cd(II)/225 mg/g/NR

pH working window [52] 4e7 Selectivity tested against Ni2þ, Cu2þ, Co2þ, Cd2þ, Ca2þ Mg2þ Pb2þ HNO3 þ thiourea as regeneration agent. Capacity loss is observed after the second cycle. NR [53] Selectivity over ca, mg, [54] Ni, Mn, Zn, Fe, Pb Regenerated with ethylenediamine solution. Capacity loss over cycling

HKUST-SO3

Single-element solution

Coordination to sulfonic motifs

Cd(II)/89 mg/g/ 10 min

UiO-66-NH2

Single element solution

Electrostatic interaction with amine groups

Cd(II)/178 mg/g/120 min

Al-MIL-53

Single element solution

Not reported

Pb(II)/100 mg/g

LN-BTC

Single-element solution

Not reported

Pb(II)/5 mg/g

MIL-101ethylenediamine

Single-element solution. Real environmental water samples, lake and river water Single-element solution

Ethyelendiamine-Pb(II) coordination

Pb(II)/80 mg/g/ 60 min

Not reported

Pb(II)/NR

pH window 4e8 [55] Selectivity proved over Na, mg, ca, Pb, Cu, Ni. Regeneration with distilled water Capacity and surface area loss at the second cycle pH window 2e7 [56] Regeneration not reported Dynamic response to [57] Pb(II) Selectivity tested [58] against many metal ions and anions found in waters sources Elution process with inorganic acid solution Elution process with [59] EDTA-2Na Capacity loss over cycling Not reported

Pd2þ/37 mg/g Pd2þ/69 mg/g Pd2þ/79 mg/g

Not reported Not reported Not reported

Ni-MOF

Weak and intermediate acidic precious ions (Pt, Au, Pd)

MOF-802 UiO-66 MOF-808

Single-element acid solution

Coordination to zirconium clusters

[60]

(Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium Compound Application conditions Adsorption mechanisms time

Fe-MIL-100Hydroquinone

E-waste leaching Gold reduction through solution obtained from polymeric the CPU (Ci ¼ 7.3 ppm) (black), wastewater (Ci ¼ 3.7 ppb) (red), and river water (Ci ¼ 120 ppb) (green)

Au(III) / Au(0)5 min

UiO-66

Acid single-element solutions (pH ¼ 1)

Coordination of MClx anions to zirconium clusters added to a partial reduction of the adsorbed species, especially in amine containing material

Simulated high level liquid waste (in 1.0 M HNO3)

Allylsulfanyl moieties affinity toward palladium coordination

Au(III) 280 mg/g Pd(II) 120 mg/g Pt(IV) 166 mg/g 10 min Au(III) 495 mg/g Pd(II) 193 mg/g Pt(III) 193 mg/g 41 mg/g

UiO-66-NH2 AS-UiO-66

Reusability

Refs.

Reusable over 5 cycles [61] Ascorbic acid activation of the adsorbent Loss of capacity over cycling recovered after activation with ascorbic acid pH working window from 1 to 11 Selectivity proved over [62] interfering metal ions including Co(II), Ni(II), Cu(II), and Zn(II)

Slight capacity loss over [121] cycling Excellent selectivity proved on real acid multi-element conditions Side adsorption of silver and selenium during the process

UiO-66-Thiourea

Simulated wastewater

Reduction of Au(III) 300 mg/g species to Au nanoparticles 200 min through electron donor groups of thiourea moieties Au-S coordination þ AuClx species coordination to zirconium clusters

Not reported

[122]

Highly acidic radioactive metals (UO2D, Th4D, CsD, Ba2D)

Al-MIL-53@AA

Single-element acid solution

Co-SLUG-35

Uranium extraction from alkaline water solutions

Multiple carboxyl and hydroxyl groups of citrate decoration motifs are the driven mechanisms of uranium and thorium chelation Replacement of anionic species at the pores by uranyl ions

UO2þ/200 mg/g Th4þ/90 min 250 mg/g

pH range 4e10 [63] Elution 0.01M HCl, HNO3, or H2SO4

UO22þ/120 mg/g/ 4h

Complete [64] transformation of the material is observed after uranium uptake A severe capacity lost is observed between the first and second cycle, but afterward the performance of the material is maintained (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium Compound Application conditions Adsorption mechanisms time

HKUST

Uranyl solutions at pH ¼ 6

Not reported

UO22þ/800 mg/g

Zn(ADC) (4,40-BPE)0.5 Ln-MOF-76

Single-element ideal solutions Single-element acid solutions

Not reported

UO22þ/325 mg/g/180 min

Not reported

UO22þ/300 min

MIL-101-CMPO

Multielement acid solution of U(IV), Th(IV), and REE(IIII)

Carbamoylmethylphosphine oxide

UO22þ/30 mg/g/10 min

Reusability

Refs.

Regeneration or reusability has not been reported pH applicability window 5e8 pH ¼ 3 ideal adsorption acidity range Selectivity tested against Pb(II), Zn(II), Cs(I), Sr(II), Cr(III), Co(II), Ni(II) The U(VI) loaded material can be eluted by 0.1 M Na2CO3 andeasily reach a desorption percent over 90% No capacity loss after 5 cycles of elution/ adsorption Adsorption and elution process were carried out at pH ¼ 4.

[65]

[66] [67]

[68]

MIL-101(Cr)triazole-COOH

Multielement acid Carboxyl groups act as UO22þ/300 mg/g/2 h solution of U(IV), and uranyl adsorption points REE(IIII) of the material. Possible monodentate and bidentate moieties between the oxygen atoms of carboxyl groups and penta-coordinated uranyl species

MIL-101-CMPO Ship in the bottle

Multielement acid Carbamoylmethylphosphine solution of U(IV), oxide molecules REE(IIII), and transition encapsulated within the metal solutions pores act as a binding sites for uranyl species

UO22þ 4 mg/g pH ¼ $3 28 mg/g pH ¼ 4 250min

pH working window [69] from 4 to 12 Competitive sorption of coexistent ions at pH ¼ 7.0 in water (Co(II), Ni(II), Zn(II), Sr(II), La(III), Ce(III), Sm(III), Gd(III), Yb(III) 1 M Na2CO3 and 1 M HNO3 used for uranium(VI) elution. After three cycles, there is a dramatic reduction of sorption capacity (ca. 50%). Compared to the fresh material Selectivity highly over [70] Al, Cd, Co, Cu, Mn, Ni, pb, Zn, Y, Eu, Gd, Nd Slight capacity loss after the elution of adsorbed ions with 0.1M HNO3 Stability of adsorbent assessed by NMR and XRD (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium time Compound Application conditions Adsorption mechanisms

MIL-101 MIL-101-NH2 MIL-101-ED MIL-101-DETA

UiO-66-AO

Coordination environment UO22þ/20 mg/g/25 min obtained by XAS in agreement UO22þ/90 mg/g/25 min with [UO2][NH2]2O2[H2O]3 UO22þ/200 mg/g/25 min and U(VI) hydroxo species. UO22þ/350 mg/g/25 min U(VI) ions are surrounded by hydroxy and water molecules, except for the amine group on the surface of the MOFs, which add the selectivity to the material due to the uranylesoft donor bases interaction

Bohai seawater and the Amidoxime motifs 2.68 mg/g seawater containing EXAFS analyses (in seawater) extra 500 ppb uranium demonstrates that the 100 min coordination numbers and bond distances are in agreement with the average local coordination environment of uranyl containing 2 or 3 amidoximes binding in a chelating fashion

Reusability

Refs.

Selectivity highly over [71] Co, Ni, Zn, Sr, La, Nd, Sm, Gd, Y at pH from 4.5 to 5.5 Effective elution with acid solution below pH ¼ 3 30% reduction of sorption Capacity for reclaimed MIL-101-ED and MIL-101-DETA compared to the fresh material pH working window 4.5 to 6 pH working [72] window 4 to 9 100% uranium elution achieved with 0.01M HNO3 solution A slight capacity loss is observed after the first cycle

UiO-66 UiO-66-NH2

UiO-68-EPAA

UO22þ/100 mg/g/300 min

Multielement acid solution of U(IV), REE(IIII) and transition metal solutions

Not reported

Uranyl extraction from sulfuric acid media

Amide coupling using an UO22þ/125 mg/g/6 h excess of carbonyldiimidazole and 2-(ethoxy(hydroxy) phosphoryl)acetic acid (EPAA)

UO22þ/115 mg/g/300 min

pH working window [73] 4.5 to 6 Ionic strength does not affect the adsorption capacity Selectivity tested against Co, Ni, Zn, Sr, La, Nd, Sm, Gd, Y UiO66 and UiO66NH2 materials are able to adsorb REE, but in lesser extent than uranyl species Adsorption kinetics [74] ([U] ¼ 1000 mg L
1, [SO42
] ¼ 1 M, pH ¼ 2 Reusability and selectivity over other elements than sulfate have not been reported (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium Compound Application conditions Adsorption mechanisms time

MIL-101-SO3H

MOF-808-SO4

MIL-101-SO3H(Cr)

Potential radioactive contamination from nuclear waste, the radioactive 137Cs Capture of radioactive barium from aqueous nuclear waste

Not reported

Capture of radioactive barium from aqueous nuclear waste

Barium chelating group (sulfate and sulfonic acid group)

Barium chelating group (sulfate and sulfonic acid group)

þ,

Cs pH 3: 23.43 mg/g pH 6: 36.47 mg/g pH 10: 29.17 mg/g Ba2þ 131.1 mg/g

Ba2þ 70.5 mg/g

Reusability

Refs.

Not reported

[75]

Ultrahigh selectivity in the presence of high concentrations of background metal ions of Csþ, Zn2þ, Ni2þ, Co2þ, Sr2þ, La3þ, and Eu3þ Captures barium irreversibly Ultrafast kinetics with kinetic rate constant k2 of 27.77 g mg-1 min-1, which is 1e3 orders of magnitude higher than existing sorbents Captures barium irreversibly

UiO-66-SO4

Byproduct of nuclear High content of sulfate fuel fission and nuclear group, whichis a strong power reactor accidents barium-chelating group is one of the most toxic radionuclides in radioactive liquid wastes

Ba2þ 181.8 mg/g

Adsorb barium ions irreversible with high stability even after gamma irradiated

Highly acidic rare earth metals separation (REE3D, Sc3D)

CMPO@MIL-101(Cr)

Ideal single-element acid solution

Not reported, but related Eu(III) with the coordination to 14 mg/g carbamoylmethylphosphine 600 min oxide motifs anchored to the organic linkers of the material

MOF-808@EDTA

Ideal single-element acid solution þ multielement solution of 19 elements

Chelation of metal ions by ethylenediaminetetraacetic acid motifs anchored to the clusters of MOF808

La (III) Punctual analyses on other REE(III) 200 mg/g 30 min

Separation factor of 3.2 [76] and 8.5 has been achieved versus Y(III) and Zn(II) ions No reusability reported Experiments performed at pH4 and 5 No reusability test [77] Study performed at pH ¼ 4

(Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium Compound Application conditions Adsorption mechanisms time

MIL-101

Multi-element acid solutions

MIL-101-NH2

MIL-101-ED

MIL-101-EDTA

MIL-101-PMIDA

MIL-101-AA

Multi-element acid solutions Multi-element acid solutions

Chelation of the metal ions at hydroxyl, amine, ethylenediamine, DETA, diethylenetriamine, PMIDA: N-phosphonomethyl) iminodiacetic acid.

Gd(III), 14,6 mg/g Nd(III), 11,3 mg/g Sm(III), 18,5 mg/g Ce(III), 17.9 mg/g La(III), 16.5 mg/g Gd(III), 21,6 mg/g Nd(III), 23,7 mg/g Sm(III), 28,4 mg/g Ce(III), 32,1 mg/g La(III), 29,3 mg/g Gd(III), 30,3 mg/g Nd(III), 36,5 mg/g Sm(III), 54,0 mg/g Ce(III), 58,9 mg/g La(III), 66,0 mg/g Gd(III), 39,7 mg/g Nd(III), 52,8 mg/g Sm(III), 67,5 mg/g Ce(III), 68,9 mg/g La(III), 73,6 mg/g Gd(III), 37,2 mg/g Nd(III), 48,3 mg/g Sm(III), 63,9 mg/g Ce(III), 66,1 mg/g La(III), 87,7 mg/g Acrylic acid functionalized Sc(III) 90,1 mg/g material. Carboxyl and amide Nd(III) 104,6 mg/g groups of acrylic acid are Eu(III) 74,9 mg/g binding positing for string acid Gd(III) 58,3 mg/g REE species, specially, Sc(III)

Reusability

Refs.

Experiments performed [78] at pH 4 and 6. Adsorption selectivity of gadolinium tested against Al, Fe, Ni, Zn, Co Reusability test performed for gadolinium adsorption tests. MIL101 materials functionalized with amine grafting motifs show a capacity fading over cycling PMIDA motif anchored to the organic linkers through amide bridges are more stable over cycling, but also a slight capacity lost is overserved after the third run

Effective between pH 1 and 4.5

[79]

0.075-AA-0.072@MIL101

Complete adsorption test on REE serie was performed.

Regeneration was tested after stripping the REE with diluted HCl acid. The Sc(III) adsorption capacity over coexisting ions was tested, following the order: Cu(II) < Co(II) NO3 e ClO4 >BF4 >CF3SO3 e ClO4 and BF4 ions Not reported

[86]

Ag-SLUG-21

Ideal Cr2O7 solution

Cr(VI) 70 mg/g

Not reported

[87]

Cr(VI) Cr(VI) Cr(VI) Disposition of chromate anions Cr(VI) after the anionic exchange determined by single crystal X-ray diffraction Anion exchange on the pores Cr(VI) of the cationic framework Anion exchange at the pores Cr(VI) of the material modifying the luminescence properties. Disposition of chromate anions after the anionic exchange determined by single crystal X-ray diffraction Anion exchange on the pores Cr(VI) of the cationic framework

Anion exchange on the pores of the cationic framework Anion exchange on the pores of the cationic framework

reported reported reported reported

[83]

Metal organic frameworks

Fe-MIL100@IL

Ideal Cr2O7 solution

NU-1000

Ideal Cr2O7 solution

MOF (Th-SCU-8)

Ideal Cr2O7 solution

UiO-66-NH2

Ideal Cr2O7 solution


UiO-66-NHþ 3 Cl

Chrome plating water sample

Anion exchange, interaction with the MOF host framework or ionic liquid components encapsulated within the material Direct hydroxyl replacement at the zirconium clusters

Cr(VI) 290 mg/g 300 min

pH 2e10 working conditions not reported reusability conditions

Cr(VI) 90 mg/g

pH working condition [89] from 1 to 7 Specific adsorption in presence of competitor anions, such as cl
, NO3, and SO4 Regeneration with acidic methanol (2.5 M HCl). Partial capacity lost in the 3rd cycle. Not reported [90]

Anionic exchange within the Cr(VI) Not reported 10 min pores of cationic framework Direct hydroxyl replacement Cr(VI) 33 mg/g 100 min at the zirconium clusters Electrostatic interaction with amine groups

Direct hydroxyl replacement at the zirconium clusters Chlorine exchange by

Cr(VI) 300 mg/g 15 min

[88]

Working pH window [91] 2e8 Effect of interfering anions, cl
, NO3, SO4, CO3, HCO3, PO3, F
; detrimental in case of highly coordinated anions. Working pH window [43] 2e8. Effect of interfering (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium time Compound Application conditions Adsorption mechanisms

chromate at NH3þCl
motifs of the organic linkers

Fe-Gallic

Ideal Cr2O7 solution

MOR-2

Ideal CrO4 and Cr2O7 solutions. Chrome platting solutions and potable water real samples

MOF-867

Ideal Cr2O7 solution

ZJU-101

Ideal Cr2O7 solution

Reusability

Refs.

anions, cl
, NO3, SO4, CO3, HCO3, PO3, F
SanddMOF Cr(VI) andorption and its Cr(VI) 1600 mg/g Working pH window [124] reduction to Cr(III) due 3-10. Effect of to the presence of hydroxyl interfering ions: groups Mn(II), Ni(II), Cu(II), Zn(II), Cd(II), Ba(II), Cr(III) Cr(VI) coordination to the Cr2O7 193.7 mg/g and Working pH window [127] Zr6 hexanuclear clusters. CrO4 118.3 mg/g < 10 min 3-8. Effect of interdering ions: proved in real water samples and plating industry solutions that contains many interfering ions. Cr(VI) chemisorption at Zr(IV) Cr(VI) 50mg/g - 150 min Not reported hexanuclear clusters. Cr(VI) chemisorption at Zr(IV) Cr(VI) 250 mg/g - < 10 min Effect of interfering [125] ions: Cl
, Br
, hexanuclear clusters þ NO3
, SO42
, I
anionic exchange process þ
and F
N eCH3Cl groups attached to the organic linkers

UiO-66-OH

Ideal Cr2O7 solution

UiO-66-(OH)2

Ideal Cr2O7 solution

Interactions and partial reduction from Cr(VI) to Cr(III) through a deprotonation of hydroxyl groups within the organic linkers. Interactions and complete reduction from Cr(VI) to Cr(III) through a deprotonation of hydroxyl groups within the organic linkers.

Cr(VI) 20 mg/g 50 min

Target metal capacity (mg/g) kynetics

Cr(VI) 60 mg/g 33 h

Working pH window 3e7 Effect on ionic strength on the adsorption No reusability test reported

[92], [123]

Reusability

Refs.

ANIONS e NEUTRAL SPECIES (As(VI)eAs(III))

Compound

Application media

Adsorption point (mechanism)

Al-MIL-53

Ideal arsenate water solutions

Fe-BTC gel

Ideal arsenate solutions

Hydrogen bonding interaction As(V)O4 100 mg/g 25 h proposed between arsenate and hydroxyl groups within inorganic chains of MIL53 material. Not reported As(V)O4 12 mg/g 10 min

Fe-MIL-53

Ideal arsenate solutions

Not reported

As(V)O4 20 mg/g 150 min

Capacity fading for [93] high coordinating anionic species, such as F-or PO43pH applicability [94] window between 2 and 10 pH applicability [95] window limited to 3 to 9. Capacity loss is observed at pH 9 and 11 (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium time Compound Application conditions Adsorption mechanisms

Reusability

Refs.

Fe-MIL-100

Ideal As(V) solutions

Not reported

[96]

In-AUBM-1

Ideal As(V) solutions

MIL-88-NH2-B

Ideal As(V) solutions

Adsorption mechanism studied As(V)O4 100 mg/g by Mossbauer’s spectroscopy. Inorganic trinuclear clusters interaction with arsenate anions Not reported As(V)O4 100 mg/g 5min

Adsorption mechanism studied As(V)O4 120 mg/G 5 min by multitechnique approach, including luminescence, which indicates that both amine groups and inorganic clusters interaction with arsenate groups exist after adsorption

Stability window [97] pH ¼ 1 to 12 Operation window ¼ pH 6e10 Activation with 7M NaCl solution. Capacity loss after three cycles Luminescence signal [98] modification was studied over the presence of Mg2þ, Al3þ, Bi3þ, Agþ, Cu2þ, Co2þ, Zn2þ, Ni2þ, Mn2þ, Pb2þand Cr3þ) and sodium salts (Ac,SO42, HCO3, CO32, NO3, Cl, Br, PO43 and AsO33.

ZIF-8

As(III) ideal solutions

Hierarchical ZIF-8 ZIF-8 nanoparticles

As(V) ideal solutions As(V) and As(III) ideal solutions

ZIF-8 with different morphologies

As(III) ideal solutions

MOF-808

As(V) ideal solutions

UiO-66

As(V) ideal solutions

Surface interaction between As(III) 90 mg/g ZIF-8 particles and As(III) species Interaction with ZnN4 centers As(V)O4 90 mg/g Direct coordinative ZneOeAs As(V)/As(III) 30 mg/g and NeOe As bonds (As(III)) 50 mg/g (As(V)) determined by IR and XPS 3h

Not reported

[99]

Not reported [100] Optimum pH ¼ 8 for [101] as(III) and pH ¼ 7 for As(V). Stability of ZIF8 compromised below pH 7. Adsorption decay is observed in the presence of highly coordinative competing anions (PO4, SO4) Direct coordinative ZneOeAs As(III) 100 e 1250 mg/g 10 h Adsorption decay is [102] and NeOe As bonds observed in determined by IR and XPS. the presence of highly coordinative competing anions (PO4, SO4) [39] Direct coordination to the As(V) 25 mg/g 50 min Capacity lost after 5 zirconium clusters cycles. Regeneration of the materials with 0.5M Na2SO4 solution Direct coordination to the As(V) 300 mg/g (pH ¼ 1) pH working window [44] zirconium clusters 120 mg/g (pH ¼ 7) 1e11 (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium time Compound Application conditions Adsorption mechanisms

UiO-66-HCl, UiO-66(SH)2 and UiO-66(SH)2eHCl

As(III) ideal solutions

UiO-66-HCl and UiO-66

As(V) ideal solutions

Direct coordination to the zirconium clusters þ interaction with thiol groups Direct coordination to the zirconium clusters

Reusability

Refs.

As(III) 15 mg/g As(III) 20 mg/g As(III) 40 mg/g As(III)

Not reported

[103]

As(V) 140 mg/g 110 mg/g

Not reported

Target metal capacity (mg/g) kinetics

Reusability

Highly coordinating Anions (PO4, SO4, FL, CNL, SeO4, SeO3, BO4)

Compound

Application media

Adsorption point (mechanism)

UiO-66 UiO-66-NH2

Synthetic human urine Synthetic human urine

Direct coordination to PO4 100 mg/g 30 min zirconium hexanuclear clusters PO4 120 mg/g 30 min

Aluminum fumarate

Fluorine synthetic solution. Real ground water

Hydroxylefluorine exchange

[{Cu(L1)2(Cl)}$ Cl$(H2O)4] (L1 ¼ N,N0 -bis(3-pyridyl)terephthalamide)

Acetonitrile solution containing [(Bu)4NF]

Anion exchange of F- /12 min noncoordinated anions trapped within a cationic framework. to exchange the encaged chloride water cluster [Cl(H2O)4]
with the partially hydrated fluoride [F(H2O)4]
. determined by single crystal X-ray diffraction

F- 600 mg/g 2 h

Refs.

Regenerated with [104] NaCl concentrated solution. Capacity fading to 85% of initial value after 5 cycles Elution processes with [105] sulfate, carbonate, Bi-carbonate, chloride solutions NR [106]

MIL-88-A

Ideal fluoride water solutions.

Anion exchange at the inorganic clusters

F-/5 min

MIL-53(Fe) MIL-53(Cr) CAU-6 UiO-66(Hf) ZIF-7 ZIF-8 ZIF-9

Ideal fluoride water solutions. 40 mg/g

Anion exchange at the inorganic clusters

UiO-66(Zr)

Ideal fluoride solution

Anion exchange at the inorganic clusters

F16.96 mg/g 10.3 mg/g 24.22 mg/g 33.35 mg/g 2.57 mg/g 0.90 mg/g 1.70 mg/g F- 40 mg/g 100 min

UiO-66 eNH2(Zr)

Ideal fluoride solution

Anion exchange at the inorganic clusters

F- 60 mg/g 5 min

Polymerization within the MIL101 porous cages. Interaction between N-Methyl-D-glucamine and borate species

B(OH)4 species 2 mg/g

MIL-101- N-Methyl-D- Ideal fluoride solution glucamine

Efficiency test against competing anions Cl
, NO3, HCO3, SO4, PO4. Not reusability reports NR

[107]

[108]

Efficiency test against competing anions Cl
, NO3, CO3, SO4. Structure is collapsed under 5 mmol fluoride concentration Cl
, Br
and surfactants [108] as competitor species has been tested. Not reusability test done. Efficiency against [109] equimolar foreign anions (NO3, CO3, SO4, PO4, Cl, Br) proved. Reusability over four without capacity fading. (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium Compound Application conditions Adsorption mechanisms time

Reusability

Refs.

SeO3, SeO4

Hf-MOF-808 NU-1000

Bi-CAU-17

[110] Ideal selenate and selenite Coordination to inorganic SeO4, SeO3 125 mg/g (SeO3) Not reported solutions Hf hexanuclear clusters 120 mg/g (SeO4) Not reported [111] Ideal selenate and selenite SeO4, SeO3 90 (SeO3) 80 through hydroxyl groups (SeO4) solutions displacement Coordination to inorganic Zr hexanuclear clusters. Experimentally proved by Pair distribution function and X-ray adsorption. Ideal selenate and selenite A computer simulation using SeO3 250 (SeO3) 25 (SeO4) Capacity fading over [112] solutions finite difference method highly coordinative near-edge structure anionic species, such (FDMNES) method under as carbonate, chloride self-consistent field (SCF) or phosphate species. mode to identify the possible Se local environment. Coordination to the clusters. Selenite species are able to displace organic linkers of CAU-17

Weekly coordinating Anions (TcO4L(ReO4L), ClO4L) ClO4L

ASC MOFs

Ideal ClO4 solution

Coordination with copper ClO4150 mg/g metal centers of the 20 min coordination polymer Interaction with aminosulfonic acid groups anchored to the copper metal centers

Not reported pH applicability window proved between 1 and 10.

[113]

Anionic exchange by sulfonate ClO4
[Cu2(bds)(bpy)2]$2H2O Ideal 0.1 M NaNO3 MnO4 [1], [Cu4(bds)2(azpy)4]$ Or NaClO4 or KMnO4 anions of the structures solutions NO3
6H2O Not reported Ideal solutions Cationic coordination polymer. ClO4
[Cu(p-bbiteb)2 (-m-Cl-[Cl$(H2O)2 Cl
anionic exchange for the NO3
3 Cu2(p-bbiteb)4])]120 min target anions take place Cl4$MeCN$(H2O)8 Capacity not reported during the process. ClO4
SLUG-21 Ideal solutions 1,2-ethanedisulfonic Not reported Acid exchange by ClO4
Ag-SCU-101 Ideal solutions Structure after adsorption ReO4
solved by single crystal X-ray 240 mg/g 30min diffraction. Anions are selectively trapped within concrete pores of the structure, nearby the cationic metal center; but not covalently connected to them. Electrostatic potential distribution of the partial framework calculated for the surrounding host adsorption position of ReO4 species Ag-SBN

Ideal solutions

Solid to solid recrystallization ReO4
of the metal-bipyridine chains 650 mg/g within the crystal structure to 25 min adapt the adsorption positions of ReO4 species after their uptake

Not reported

[114]

Not reported

[115]

Not reported

[116]

Adsorption capacity [42] tested against competitors such as NO3, CO3, SO4, PO4, NO3 Efficient even in the presence of sulfate concentration relevant for real application conditions Performance maintained after the exposure of the material to gamma radiation pH applicability [117] window between 3 and 9 Adsorption is effective in the presence of high nitrate concentrations (Continued)

Table 3.2 Summary of metal ions adsorption capacity, kinetics, reusability, and adsorption mechanisms.dcont’d Target metal/capacity (mg/g)/equilibrium time Compound Application conditions Adsorption mechanisms

Ag-SCU-100

UiO66-NH3þCl


Ideal solutions

Direct coordination of ReO
4 to the silver metal ions

ReO4
375 mg/g 30 min

Replacement of chlorine ions at the NH3þCl
motifs for ReO4 species located by electronic density map exploration obtained from synchrotron based X-ray diffraction data.

ReO4
159 mg/g 24 h

Reusability

Refs.

pH applicability [118] window between 1 and 13 Stability confirmed after gamma and beta radiation. High rates of adsorption efficiency retention in presence of high concentrations of NO3 and SO4 salts Selectivity over ReO4 [119] diminish in presence of phosphate anions

Porous, lightweight, metal organic materials: environment sustainability

93

sulfur containing organic moieties. Fig. 3.10 summarizes maximum adsorption capacities reported for MOFs over mercury, cadmium, and lead from aqueous solutions. Sulfur-containing moieties are known to have high affinity to mercury chelation. Indeed, zirconium terephthalate UiO-66, one of the iconic water-stable MOFs, has been widely pre and post-synthetically modified, both at organic and inorganic building units, incorporating this type of soft chelating motifs to its framework. Terephthalate pillars of UiO-66 have been functionalized with thiol, thio-urea, isocyanate, thiocyanate, amine, and amine motifs able to capture in different extents mercury from aqueous solution. The density of these anchoring points within the porous framework is not the unique criteria influencing the mercury adsorption, since, for example, UiO-66-(NHC(S)NHMe)0.73, which combines both amine and thiol groups in a branch-like mercury chelating trap, is much more effective than UiO66-(SH)2 compound to capture mercury, even when the density of adsorption point is much more lower. Inorganic cluster modifications through solvent-assisted ligand exchange have demonstrated to be also an effective approach to incorporate Hg(II) adsorption acetate thiol motifs to UiO-66 structure. Within the works exploring the cluster decoration approach as potential path to incorporate metal chelators within MOFs highlights the mesoporous MOF-808 postfunctionalization with the tetrapodal ethylenediaminetetraacetic acid (EDTA), a commercially applied heavy metal chelator. Indeed, MOF808@EDTA shows a high capacity and affinity to capture from soft Hg(II), Cd(II), Pb(II) to highly acidic Cr(III), REE(III) metal ions. 1000

Hg(II)

Pb(II)

Cd(II)

ZIF8

Zn-PCN-101 Zn-TMU16-NH2

FJI-12 LNMOF-263 Zn(hip)(L)

Zn-PCN-101 FJI-9

HKUST-SO3 Ni(3-bpp)2 (NSC)2

HKUST HKUST -dithioglycol

Ln(BTC)

Cu4 (Methionine)2 Ca2Cu4(methionine)2

Cr-MIL101 alkenyl -thiol Cr-MIL101-NH2 Cr-MIL101 - thymine AI-MIL53-MNP

DUT-67

MOF808 - EDTA

UiO66-NCS UiO66-NCO UiO66-NHC(S)NPh UiO66-MAA

400

UiO66-NHC(S)NHMe

600

UiO66 UiO66-NH2 UiO66-(SH)2

Adsorption capacity (mg/g)

800

200

0

Figure 3.10 Maximum adsorption capacities reported for Hg(II), Cd(II), and Pb(II) metal ions for MOF and MOF-modified adsorbent materials.

Ainara Valverde et al.

94

Other water-stable MOFs based on sulfur containing linkers, such as, DUT-67, FJ-9, and FJ-12 have been revealed as efficient mercury adsorbents. It is worthy to mention that even having a lower adsorption capacity, thymine modified Cr-MIL-101 material is highly selective to capture mercury even in real water samples, which further confirms the feasibility of MOFs application in real conditions. One of the most interesting examples of mercury adsorbents are the copper Bio-MOFs based on methionine natural amino acids. Copper and calciumecopper mixed Bio-MOFs structures arrange the metal ions and carboxylate groups of the methionine linker in a three-dimensional structure showing 1D channels toward the S-methyl thioester side chain is pointing. Besides their record adsorption capacity, water stability (even being a divalent metal based coordination polymer), and mercury selectivity and reusability; copper Bio-MOFs are obtained as single crystal that enables tracking the mercury speciation and positioning within the MOF structure. Single crystal X-ray diffraction experiments reveals that HgCl2 species are trapped within the pores or the material forming HgCl2S2 adducts with the methyl thioether side chains of the methionine linkers (Fig. 3.11).

(a)

(b)

S H C N O S CI Ca Cu Hg

Figure 3.11 Crystal structure of CaCu-BioMOF and local coordination environment of HgCl species at different stages of the adsorption process (a) 3HgCl2 and (b) 5HgCl2.

Porous, lightweight, metal organic materials: environment sustainability

95

Independently on cluster or organic linkers anchored functionalities, the sulfur containing moieties are quite effective for mercury adsorption, but they work in a less extent for cadmium or lead ions. Indeed, the presence of adsorption points with higher pKa values than the ones of sulfur moieties, such as amine or carboxylate groups within EDTA, hydroxyl groups within the inorganic chains of Al-MIL-53 structure, or sulfonyl groups of HKUST-SO3 compound, favor in more extent the adsorption of Cd(II) and Pb(II) species. Among the cationic MOFs, AMOF anionic coordination polymers compensate the net negative charge of the network with cationic organic solvents trapped within their pores. As cationic exchange polymeric resins, AMOFs can interchange the organic cation for mercury, cadmium, or lead ions. Nevertheless, the kinetics of this cationic exchange process are quite slow, with equilibrium adsorption and desorption times up to 1 day (Fig. 3.12). 3.4.1.2 Platinum group soft metal ions recovery from different media Platinum group metals are soft acid ions, but they are stable as PtCl4, PdCl6, and AuCl4 anionic species at HCl acidic conditions obtained after the acid leaching of electronic scratch or other end life components, a common step to proceed to extract and recover these economically and technologically relevant ions. Fig. 3.13 summarizes the adsorption capacity of different MOFs over Pd(II), Pt(II), and Au(III). The first study tackling to recover these highly valuable metal species was developed from ideal single-element solutions with UiO-66 and UiO-66-NH2 materials (Fig. 3.13). The high capacity (up to 450 mg/g for Au, and near 200 mg/g for Pd and Pt) and fast kinetics ( Phosphate > Acetate > Citrate > Tartrate > Hydrogen > Carbonate > Chromate z Chloride > Nitrate > Chlorate. Up to date it has been difficult to disrupt this selectivity rule with technologies based on polymer chemistry, but the crystal chermistry versatility of MOFs and coordination polymers has offered the possibility to circumvent the general tendency described by Hofmeister rule. Before the discovery of the MOFs understood as were defined in 1999 [41], several studies were reported on anionic exchange in positively charged metal-ligand porous frameworks encapsulating anionic solvents or inorganic anions within their pores. Usually, the cationic inorganic-organic compounds are formed by the combination of neutral organic linkers (e.g., pyridine type) and transition metal ions. Depending on the catione linker molar ratio, and their connectivity, a broad scope of coordination polymers showing from one-dimensional to three-dimensional noninterpenetrated and interpenetrated frameworks have been reported so far. As a consequence of their crystal-chemistry versatility, cationic coordination polymers have shown a wide variety of framework to anion interactions, ranging from weak hydrogen bonding and nonspecific van der Waal’s forces to coordination to Lewis acid metal centers. In addition to the wide variety of interaction between the oxyanions and coordination polymer’s structure, the functionalization of the framework with strong directing hydrogen bonding groups complementary to the included anions, or the possibility to tune the rigidity or chemistry of the pores, allow disrupting the Hofmeister selectivity rule, to obtain specific anion uptakes, and henceforth, anion separation paths. As a general rule, transition metal cationic polymers show slow anion exchange kinetics, with equilibrium times above 1 h in the most favorable cases, and more commonly up to one or several days. Anion exchange kinetics in metalecoordination polymers depends also on the anion species migration paths to the final adsorption points. Frameworks with wider pore windows, as well as a labile framework that can be adapted to the migration of oxyanions, favors the anion exchange kinetics. In view of some of the chemical similarities between the cationic coordination polymers and permanent porous metal organic frameworks, the potential of MOFs as inorganic anion adsorbents has been widely explored aiming the removal of environmentally hazardous anions from polluted waters. The presence of replaceable anionic species in the inorganic clusters

104

Ainara Valverde et al.

offers the opportunity to explore anion exchange processes within these materials. Indeed, MOFs have shown outstanding results in term of capacities, kinetics, and selectivity for capture of many of the anionic pollutants studied before for cationic metal-organic polymers. Giving a step further, some water-stable MOFs have been chemically modified to mimic the anion exchange strategies reported in cationic inorganic polymers. To this end, positively charged organic groups have been included within the organic linkers, the net positive charge of the network being compensated by anionic species located at the pores, which lead to an increase of the capacity and selectivity of the modified frameworks. MOFs and cationic coordination polymers’ performance over the removal of phosphate, fluoride, arsenate, arsenite, chromate, selenate, selenite, perchlorate, pertechnetate, or permanganate has been studied in different scenarios, ranging from simulated mono-anionic solutions to real polluted water samples, such as chromium containing solutions obtained from plating industry. Nevertheless, it is necessary to consider the solution chemistry beyond the application. The speciation of oxyanions in solutions makes the difference between the mechanisms of exchange or adsorption observed in coordination polymers and MOFs. Therefore, depending on the pH and eH speciation of the above mentioned oxyanions, we consider necessary to structure the following section of this chapter in the following blocks: (i) Hexavalent chromium oxyanions (CrO42
and Cr2O72
) adsorption and reduction to cationic chromium trivalent species. (ii) Arsenate (As(V)) and arsenite (As(III)) oxyanions, and arsenite (As(III)) neutral species adsorption. (iii) Strong and weak coordinating inorganic anions mainly present as X
and HyXO4
n. 3.4.2.1 Chromium oxyanions capture by metal organic frameworks and related materials Chromium is usually present as CrO4
2, Cr2O72
, and HCrO4
highly mobile and soluble hexavalent oxyanion, or as less soluble trivalent Cr3þ oxo-aqua species (mainly stable at acidic conditions). Hexavalent chromium and its compounds are carcinogens, since their oxidizing capacity is able to provoke damages in the DNA during its reversible oxidation/reduction to highly reactive intermediate Cr(V) and Cr(IV) species. Chromium is mainly anthropogenically incorporated within the environment through various industrial processes, such as metal plating, leather tanning, or cement production (Fig. 3.18).

Porous, lightweight, metal organic materials: environment sustainability

(a)

(b) 1200 1.2

–2

HCrO4

H3AsO4°

–2

Cr2O7

800

0.8

H2AsO4–

–2

Cr(OH)3(s)

0.0

CrO4

400

2–

HASO4

3–

AsO4 H3AsO3°

0

lv

O2 ed

–0.4

Eh (mV)

Reducti on b

y

MnO2 disso

Eh, volts

Cr+3

0.4

–400

Cr6+ stable

H2AsO3

3–

AsO3

Cr3+ stable

-0.8

Source: Bodek2

0

2

HASO3

–800 4

6

8

10

12

14

0

2

4

pH

6

8

10

2–

12

14

pH

Figure 3.18 Chromium and arsenic speciation eH-pH diagrams.

105

106

Ainara Valverde et al.

Among the current solutions to stabilize and safely dispose Cr(VI) anions, reduction and precipitation methods are the most promising ones, since they finally immobilize the chromate anions as Cr(III) insoluble oxides/hydroxides. Nevertheless, in many scenarios Cr(VI) to Cr(III) chemical reduction process is ineffective in lowering the chromium levels in water below the legal limits. As alternatives, adsorption or ion exchange represent a promising, clean, economic, and reusable means to this end. MOF adsorbents have overcome the capacity values reported for classic adsorbents, anion exchangers, or polymeric resins to retain chromium species. Fig. 3.19 summarizes the adsorption capacity and equilibrium times of cationic coordination polymers and metal organic frameworks over chromate anions. Cationic coordination polymers for chromate immobilization

Cationic metal organic frameworks are usually formed from a combination of neutral or anionic organic linkers and metal cations, giving rise to a net cationic metaleorganic network, the charge of which is compensated with anionic counter-anions. Anionic exchange in cationic coordination polymers, and its potential in oxyanion water remediation, was known much earlier than the discovery of MOFs. Indeed, depending on the assembly of metal and neutral linkers, cationic coordination polymers with three-dimensional (1$ClO4, ABT$2ClO4), two-dimensional (FIR53, FIR-54), and one-dimensional (SLUG) crystal structures have been obtained and studied as chromate exchangers. Bimetallic 3D 1$ClO4-MOF is formed by Dy2(COO)6(H2O)4 and [Zn(BPDC)3] units connected through an organic linker containing carboxylate and pyridine metal binding motifs, and finally forming a 1D porous cationic structure where the ClO4 anions are located at the pores of the frameworks as counteranions. Chromate exchange in 1$ClO4-MOF is fast, selective, and moderately effective (70 mg/g) in comparison with other reported compounds. ABD$2ClO4 cationic compound exhibits interesting response over chromate uptake. Its structure is assembled from nanoscale coordination cages generating 1D channels with 7.8 Å diameter, where ClO4 anions are located. Interestingly, the exchange process between ClO4 and Cr2O7 has been followed by single crystal X-ray diffraction, identifying the exact position of chromate anions within the silver tetrazolate porous cages after the exchange process.

350

300

250

200

150

100

50

0

NU-1000 UiO-66

UiO66-NH2 UiO-66-NH3+CI–

Cr(VI)

UiO66-(OH) 2

Equilibrium time

UiO66-OH

ZJU-101 MOF-867 Fe-MIL-100-IL Fe-gallic MOR-2 (Cr2O7) MOR-2 (CrO4) HKUST 1.CIO 4

ABT.2CIO4 1.SO4 Ag-triazole-1 Ag-triazole-2 SLUG-21 Co-SLUG-25 ZnCo-SLUG-25 MONT-1 FIR-54 FIR-53 Cd-MOFs 3000

2500

2000

1500

1000

500

Equilibrium time (min) Porous, lightweight, metal organic materials: environment sustainability

0

Figure 3.19 Maximum adsorption capacities and equilibrium times reported of chromate anions adsorption for MOF and MOF modified adsorbent materials.

Cr(VI) adsorption capacity (mg/g)

107

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Zn and nickel tris(4-(1H-imidazole-1-yl)phenyl)amine) 2D FIR-53, and 3D FIR-54, and 1$SO4 cationic networks possess 1D pores with large apertures. The counteranions are located at the surface of the channels compensating the cationic charge of the framework. Depending on the pore aperture of FIR-53, FIR-54, or 1$SO4, a moderate to fast and effective chromate uptake has been reported, with capacities ranging from 75 to 225 mg/g. One of the main advantages of FIR-53 and 1$SO4 compounds is that the anion exchange has been followed by single crystal X-ray diffraction, and the location of chromate species after the exchange process has been accurately determined. A new paradigm for anion exchange processes was reported by L. Zhu et al. [42] for cationic coordination polymers formed by silver (SLUG-21), zinc, and cobalt (SLUG-25) 1D 4,40 - bipyridine cationic compounds. L. Zhu et al. [42] revealed that the packing of the metal ligand chains can be rearranged accommodating different anionic species. The process can be defined as a solid to solid recrystallization driven by the anion exchange process. Similar strategy was observed to trap chromate on MONT-1 {[Ag(m3-abtz)]$(NO3)$(0.125H2O)}n (1-(4-aminobenzyl)-1,2,4-triazole (abtz)) compound. Robust metal organic frameworks for chromate capture

Initial MOF adsorption experiments over chromate anions were performed with Zr-NU-1000, Zr-UiO-66, and Fe-MIL-100 iconic materials. Hexanuclear zirconium clusters of UiO-66 and NU-1000 are able to undergo an anionic exchange process between the hydroxyl groups of the Zr6 units and the chromate anions. Indeed, the superior chromate adsorption capacity and kinetics of NU-1000, in comparison to UiO-66, UiO-67 and UiO-68, was ascribed to the presence of ordered defective positions, not occupied by carboxyl groups, within the inorganic clusters, combined with large pore window apertures. Local structural models corroborating the coordination mechanism of the chromate species to the Zr6 clusters have been reported for similar anionic selenate and selenite species in NU-1000. Furthermore, local structural models simulated through molecular modeling also support this chromate chemisorption mechanism in Zr-based MOR-1 adsorbent (Fig. 3.20). Mimicking the anion exchange mechanics reported in cation coordination polymers, amine groups within UiO66-NH2 and UiO68-NH2 frameworks have been protonated after the exposure of the initial materials to hydrochloric acid solutions, giving rise to NH3þCl
pairs within their

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Figure 3.20 Chromate anions local structure after their binding to Zr6 hexanuclear clusters of MOF-1.

structures. Indeed, the UiO66- NH3þCl
compound shows an adsorption capacity nearly sixfold higher (250 mg/g) than the nonprotonated UiO66NH2 material, since, added to the chromate anions affinity to bind zirconium hexanuclear units, they can also exchange chloride anions of the NH3þCl
pairs located at the organic linkers. Similar approach has been followed for ZJU-101 material, but applying a post-synthesis protonation of the zirconium-2,20 -bipyridine-5,50 -dicarboxylate linkers (MOF-857). MOF-857 is an expanded structure of UiO-66 material, analog to UiO67 framework. Z. Wang et al. [43]. confirmed that UiO-66-(OH)2 combined the chromate coordination ability of zirconium hexanuclear clusters with the Cr(VI) to Cr(III) reduction capacity of hydroxyl electron donor groups at the organic linkers. The dual adsorption-reduction mechanism not only improves the chromate adsorption capacity of the bare UiO-66 material but also immobilizes chromium ions as a mixture of hexavalent and trivalent less toxic species. It is worthy to mention that chromate anionic adsorption in zirconium oxides supports, or zeolite materials, generates a chromium speciation that goes from hexavalent to trivalent ions, stabilizing even highly reactive and unstable pentavalent species through the generation of Zr-O-Cr bridges. The well-known oxidation-reduction lability of chromium species anchored to zirconium oxide surfaces probably can be replicated at

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Zr-MOF type materials. Indeed, it is foreseen that equilibrium between Cr(VI), Cr(V), and Cr(III) ions will be stabilized within the Zr-based MOF frameworks, depending on the reduction capacity of the groups surrounding the adsorption positions. In that respect, EPR and UV-VIs spectroscopy studies developed by G.-Sainz at al. point that chromate adsorption in UiO-66-R (-H, NH2, -(OH)2) type materials goes beyond a simple chemisorption process, since Cr(VI) adsortion is coupled to its reduction to Cr(III) through Cr(V) intermediate species, and the posterior clustering of Cr(III) ions within the UiO-66-R (-H, NH2, -(OH)2) frameworks.[123] 3.4.2.2 Arsenate and arsenite species Arsenic is a highly toxic, long-term, persistent metalloid released to the media by the combustion of fossil fuels, metallurgy, and mining industries. Natural arsenic surface or underground water pollution also occurs because of its presence in local bedrocks worldwide. Indeed many countries present natural arsenic concentration in water above the 10 mg/L legal limit established by World Health Organization (WHO). Depending on the oxidation/reduction conditions and pH of the media, arsenic is present as oxyanions based on As(V) or As(III) neutral or negative species. Due to the neutral character of As(III) species, they are more difficult to remove through adsorption processes than arsenate anions, being the chemical versatility of MOFs and advantage to tackle this challenge. The main studies carried out for arsenate adsorption have been based on iconic zirconium (UiO-66, MOF-808), chromium and iron (MIL-101, MIL-53, MIL-88), and zinc (ZIF-8) iconic MOFs. Punctual but very promising studies have been developed for As(III) adsorption with ZIF-8, UiO66SH2 analog, and Cr-MIL-101 materials (Fig. 3.21). As reported for chromate, arsenate adsorption in zirconium MOF materials is based on the direct coordination of arsenate groups to the Zr6 oxo-clusters through the replacement of one hydroxyl group. This hypothesis was further confirmed by C. O. Audu et al. [44] correlating the arsenate adsorption capacity of UiO-66 materials with the linker defect degree induced in the crystal structure through HCl modulation. A pH dependence on adsorption capacity was reported by C. Wang et al. [39] for UiO-66, with a maximum capacity values of 300 mg/g reached at acidic conditions (pH ¼ 2). Giving a step further, MOF-808 or NU-1000 has been proved to be effective to diminish the concentration of arsenate species below the legal limits in real water samples.

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350 As(V)

As(III)

CoFe2O4@Cr-MIL101

CoFe2O4@Cr-MIL101

Fe3O4@Cr-MIL101

ZIF-8

MIL88-NH2

Fe-MIL100

Fe-MIL53

AI-MIL53

AUBM-1

MOF808

UiO67

150

UIO66

200

UIO66-(SH)2

250

UIO66 – pH1

As(V)-As(III) adsorption capacity (mg/g)

300

100

50

0

Figure 3.21 Arsenate (As(V)) anions and arsenite species (As(III)) adsorption by MOF and MOF modified materials.

Iron, chromium, and aluminum MOFs based on trimeric inorganic clusters, such as MIL-100 and MIL-101 materials, have shown similar adsorption capacities than Zr-based MOF, exhibiting near 100 mg/g adsorption capacities and fast kinetics below 1 h of equilibrium times. Post-adsorption characterization of trivalent MOFs by infrared and X-ray photoelectron spectroscopies suggest that the arsenate adsorption is related to the hydroxyl groups exchange at the trimeric clusters. This hypothesis is further confirmed by M€ ossbauer spectroscopy spectral signals associated for the iron species within Fe-MIL-100 material after arsenate loading, showing a larger quadrupole splitting for one of the iron cation, over the three existing in the trimeric units. The authors speculate that this experimental evidence is ascribed to the displacement of a coordinated water molecule by an arsenate group at one of the octahedral positions with the Fe-trimeric clusters (Fig. 3.22). Regarding the adsorption of neutral (pH < 7) or negative (pH > 7) arsenite species, only three MOF materials have been used to this end. UiO-66-(SH)2 thiol functionalized analog shows that the inclusion of basic weak groups induce an increase of the affinity toward neutral arsenite species

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Fe3(OH)2(H2O)(–COO)6 HAsO4

Figure 3.22 Illustration of arsenate interaction with trimeric inorganic units of FeMIL100 material.

up to 15 mg/g of maximum capacity. Magnetite@FeMIL-101 and Magnetite@CrMIL-101 composite show a quite good affinity and capacity toward As(III) uptake, but it is mainly attributed to the arsenic immobilization in the magnetite nanoparticles, since the Fe-MIL-101 capacity is negligible and Cr-MIL-101 shows also a low arsenite adsorption capacity (20 mg/g). ZIF-8 material has been revealed as the most effective adsorbent for neutral and negative arsenite and arsenate species at neutral and basic pH values. Given the pore aperture of ZIF-8 cages (3.5 Å), in comparison to the arsenite and arsenate anions (w7 Å), it is foreseen an arsenic uptake mechanism at the surface of the ZIF-8 nanoparticles through Zn-O-As covalent bridges or N$$$H-O-As and N-H$$$OAs hydrogen bonding interactions. In general, terms, the MOF’s selectivity toward arsenic species decrease in the presence of phosphate and carbonate highly coordinative competing anionic species. Regeneration test indicates that the studied MOFs are mostly reusable without a drastic loss of adsorption capacity, with the exception of ZIF-8 material, since a partial dissolution of the material occurs during arsenic elution. 3.4.2.3 Other anionic species Besides arsenate and chromate, the adsorption or anion exchange capacity of MOFs has also been explored for other anions, such as pertechnetate,

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permanganate, chlorate, phosphate, sulfate, or fluoride. The adsorption mechanisms reported for these anions can be differentiated in function of their strong or weak coordinating character; since highly coordinating species will be stabilized covalently attached to the inorganic building blocks of MOF; while weakly coordinating anions prompt to be immobilized through weak interactions or after an anion exchange process occurring at the pores of cationic MOFs. Fig. 3.23 summarizes the adsorption capacity of the MOF porous materials studied to adsorb these oxyanions. TcO4
, ClO4
, MnO4
, CN
weak coordinative oxyanions

Tecnetium(VII) isotope is generated as a final byproduct of uranium and plutonium fission reactions. Its safe disposal is one of the major concerns, since 99Tc generates, at a long term (b emitter with a half-life of 2.13  105 years), the greatest radiation dose in the vadose zone of the waste repository. Tc is stabilized as weakly coordinative and highly mobile pertechnetate TcO4
species, which further difficult is its immobilization as solid state. 99

700

PO4

TcO4

F

CIO4

ZIF-9

ZIF-7

MIL88

CAU-6

AI-Fum

Cr-MIL53

Fe-MIL53

SLUG-21

UiO66

UIO66-NH2

SBN

UiO66-NH3+Cl–

300

SCU-101

400

SCU-100

500 ACS-MOF

Adsorption capacity (mg/g)

600

200

100

0

Figure 3.23 Adsorption capacity of cationic coordination polymers and metal organic framework materials over TcO4
, ClO4
, F
and (HxPO4)
3þx anions.

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Therefore, it could be highly desirable to eliminate TcO4
before the vitrification process of radioactive waste in order to prevent its volatilization as Tc2O7, which would require from highly capacitive, fast, and stable adsorbents able to work in caustic, high sulfate content and radioactive media. Perchlorate and cyanide anions have also a great impact on the environment and human health even at very low doses. Cyanide pollution arises mainly from precious metal mining, while perchlorate is released to water from various industrial processes. Following the same strategy applied to remove chromate anions from water, pertechnetate TcO4
species has been effectively retained applying an anionic exchange process in cationic coordination polymers (SLUG-21, SBN, SCU-101, SCU-100, ACS-MOF). SLUG-21 cationic silver-based MOF, previously studied for chromate adsorption, shows a high capacitive (w600 mg/g) and fast capture TcO4
species. Exchange mechanism is similar to chromate anions, retaining the pertechnetate species within the pores of silver-4,40 -bipyridine (Bpy) network, and nearby the silver cation. The pertechnetate record adsorption capacity of 786 mg/g has been reported for Ag(Bpy)NO3 (SBN) cationic coordination polymer. Pertechnetate capture by SBN can be understood as a recrystallization process, since during the nitrate exchange, the silver-4,40 -bipyridine metal organic chains are completely rearranged within a new crystal structure to host the TcO4
anions; as derived from the Ag(Bpy)TcO4 structure obtained from single crystal X-ray diffraction (Fig. 3.24). This coordination polymer shows a quite broad application window in terms of pH, from 3 to 9, with a high selectivity toward TcO4
in the presence of NO3. Nevertheless, the irreversibility of the anion exchange process impedes its reutilization; and in addition, an important dissolution of SBN and SLUG-1 materials over operations conditions is observed. In order to prevent the structural collapse of bipyridine-type cationic frameworks, silver and nickel cationic MOFs based on tetra-imidazole (i.e., tetrakis[4-(1-imidazolyl)phenyl]methane) linkers has been also studied for technetium removal; since tetrapodal tetrakis[4-(1-imidazolyl)phenyl] methane linkers increase the connectivity of the obtained networks, a key feature to enhance the chemical stability of MOF materials. SCU-100 [Ag2(tipm)]$2NO3$1.5H2O (SCU ¼ Soochow University, tipm ¼ tetrakis[4-(1-imidazolyl)phenyl]methane) compound has two coordination geometries for silver ions, but its connectivity with the tetradentate tipm linker generated a three-dimensional eightfold interpenetrated

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(a) ReO4–

(c)

115

(b)

NO3–

Anion exchange

ReO4–

NO3–

(d) Phase transformation

Figure 3.24 Structural rearrangement of metal-bipyridine chains in SLUG-1 material after the nitrate to techsenate anion exchange process. (Reprinted with permission from Ref. [117]. Copyright (2017) American Chemical Society.)

network with 6.9  6.9Å channels along the [001] crystallographic direction, where the nitrate anions and water molecules are located. The nitrate to pertechnetate exchange process was driven in a mono-anionic solution, taking place below 10 min. SCU-100 exhibits an excellent capacity (350 mg/g) and selectivity even in Hanford LAW Melter Off-Gas Scrubber Solution, where nitrate, nitrite, sulfate, and phosphate anions are 100 fold more concentrated than pertechnetate. As a handicap, it is worthy to mention that reusability tests have not been reported. Authors claimed that the high capacity, fast kinetics, and selectivity of SCU-100 are based on a combination of a microporous framework with the presence of weak acid silver ion open metal sites, favorable to interact with weak bases as TcO4
. The engineering of nickel cationic polymers exhibiting high connectivity, as is the case of SCU-101 [Ni2(tipm)2(C2O4)](NO3)2$2H2O), offers to meet the high capacity and fast kinetic of TcO4
exchange to the desired chemical stability and reusability for TcO4 recovery and safe disposal purposes. Combination of nickel and tipm generates a 3D cationic framework with different pore apertures were nitrate counter anions are located. The exchange process of nitrate for TcO4
is only effective for the nitrate groups located at the channels with larger apertures, which limits the exchange capacity of the material, 240 mg/g, in comparison with the above

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described materials. Although this is a drawback, SCU-101 is stable at highly acidic and basic solutions (pH range 2e14), and after exposing to ionizing radiation fields (200 kGy 60Co g irradiation or 200 kGy b irradiation (1.2 MeV)). In addition, the material can be nearly fully regenerated exchanging back the TcO4 using a 1M NaNO3 solution. UiO-66-NH3þCl
cationic MOF has been also used for ReO4
anion sequestration, showing an appreciable capacity of 160 mg/g, but slow kinetics with equilibrium times up to 24 h. Selectivity over ReO4
uptake was also severely diminished in the presence of sulfate or phosphate anions with high density charges. It is worthy to mention that ReO4
location with the UiO-66 framework was obtained through the residual map densities interpretation obtained from synchrotron radiation powder X-ray diffraction analyses. ReO4
sits within three benzene rings, along the diagonal line of the cubic cell, approximately 3 Å away from the m-O atom of the [Zr6O4(OH)4]12þ cluster; which confirms a Cl- replacement and discards a direct coordination of rhenium atoms at the zirconium clusters (Fig. 3.25). Perchlorate capture through anion exchange is also possible with SLUG21, reaching a maximum adsorption capacity of 125 mg/g. In this regard perchlorate exchange in other cationic MOFs has been also reported, but lack of adsorption isotherms and kinetic studies hinders their comparison with SLUG-1 material. Indeed, despite no reports being published for cyanide adsorption, the existence of many cationic MOFs with cyanidetype groups weakly coordinated to the metal ions suggests that they could be used also to capture these dangerous pollutants from aqueous media. PO4, SO4, and F- high coordinating oxyanions

Phosphate and sulfate pollution arises mainly from agricultural industry. In addition, there are several scenarios where the removal of sulfate anions could be desirable, such as in water solutions of spent radioactive waste, since sulfate hinders the posterior vitrification radioactive waste process needed to store them more safely at the solid state (Fig. 3.26). Phosphate and sulfate adsorption studies are limited to some zirconium frameworks, such as iconic UiO-66 and MOF-808 materials. Complete PO4 thermodynamic adsorption study has been developed for UiO-66 and UiO-66-NH2, showing quite fast kinetics with equilibrium times below 20 min, and reasonably high adsorption capacities near 160 mg/g for UiO66-NH2. Although the adsorption mechanisms have not been reported, it is expected to be similar to the arsenate or selenate chemisorption mechanism at the zirconium hexanuclear clusters reposrted for other

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b a c

Figure 3.25 Differential Fourier density map of UiO-66-NH3þ (top right) and UiO-66NH2 (top left). The ball in UiO-66-NH3þ represents Re (bottom) The position of the Re atom with respect to the basic building unit of UIO-66-NH3þ is shown. (Reprinted with permission from Ref. [119]. Copyright (2016) American Chemical Society.)

zirconium based MOFs. Indeed, the snapshot of the MOF-808
s zirconium clusters after their decoration with sulfate groups, obtained by single crystal X-ray diffraction, can be a quite accurate approximation to described the local structural model also for phosphate or arsenate anions. Given the chemical formula of the MOF-808@SO4, it can be expected a maximum sulfate adsorption capacity around 174 mg/g. Fluoride containing minerals, and some manufacturing processes using fluorine compounds, are the main origins of fluorine pollution in water.

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(a)

(b)

Figure 3.26 MOF-808@SO4 structure solved by single crystal X-ray diffraction.

Although fluorine can be removed from water using precipitation, electrodialysis, and nanofiltration, anion exchange is preferred due to its easy and low-cost operation. From a crystal-chemical point of view, the high electronegative character of fluorine ions makes them highly coordinative ions that tend to complete the coordination environment of inorganic cations, such as transition metal centers used to crystallize MOF materials. Indeed, several studies have reported the lability of trivalent metals on trimeric oxo-clusters or metal aqua-oxo chains to complete their coordination environment with different anions, depending on the synthesis conditions. Fluoride adsorption capacity has been studied for iron, chromium, and aluminum MIL-53-type materials based on terephthalate and fumarate linkers, as well as for the iconic Fe-MIL-88. The complete study of fluorine adsorption by UiO-66 and UiO-66-NH2 materials has been also performed. Reusability studies have only been reported for aluminum fumarate material, achieving a 70% of elution degree of fluorine after adsorption. The main fluoride adsorption mechanism is ascribed to the fluoride anion exchange process of the hydroxyl groups located at the inorganic clusters of the MOF materials. Assuming this, the fluoride theoretical maximum adsorption capacities can been calculated if a full replacement of hydroxyl groups per fluoride anions takes place, as reported in Table 3.3. These values are in good agreement with hydroxyl to fluorine groups replacement rated between 10% and 62%, with the exception of aluminum fumarate and MIL-88 compounds, for which the experimental capacity exceeded significantly the maximum

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Table 3.3 Theoretical and experimental fluoride adsorption values for different MOF materials. Theoretical Experimental % Exchange

Al(OH)(fum)$xH2O Al(OH)(BDC)$xH2O(CAU) Fe(OH)(BDC)$xH2O Cr(OH)(BDC)$xH2O M3IIIO(H2O)2X(BDC)3 (X ¼ Cl, F, OH) UiO-66 Zr6(OH)4O4(C6H4O4)6 UiO-66-NH2 Zr6(OH)4O4(C6H4O4NH2)6

120 102 102 102 26 e 52

555 24 17 10 40

462% 24% 14% 9.8% 154%/72%

50 mg/g

33

66%

47 mg/g

20

43%

theoretical value. For MIL-88, this experimental evidence can be explained if a partial replacement of tri-coordinated oxygen atoms located within the trinuclear units is assumed. However, for aluminum fumarate material, an additional mechanism needs to be considered to explain the experimental values.

3.5 Future perspectives and concluding remarks Among the current strategies applied to recover metal ions from aqueous solutions, metal chelators grafting to inorganic and organic components of MOFs, or the generation of positively/negatively charged metal-ligand frameworks or motifs within the MOF frameworks, are the most successful ones in terms of capacity and selectivity. In general terms, the MOFs’ key structural encoding strategies that led to the controlover capacity, selectivity, and kinetics are well established. Some of them have mimicked previous strategies applied to functionalize classic adsorbents, but in comparison, with outstanding results when using MOFs asa platform. Nevertheless, there is a room for improvement in terms of metal ions capture mechanisms elucidation, as well as accurate step-by-step determination of MOF materials chemical stability, understood as a multiple concept encompassing structural, porous, and specific functionalization motifs stability and modifications along the adsorption, desorption, activation, and reusability tests. Once established that the acidebase theory for cationic species is also effective within MOF porous frameworks, it would be interesting to

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determine the effect of the crystal framework to disturb or enhance this general rule in terms of metal ion affinity/selectivity gaining. The same concept would be also key to understand the effect of a crystal ordered environment on the Hofmeister rule disruption for anionic species. Multivariate functionalization of MOFs could be a promising strategy, proved to be highly effective for gas/drug capture and controlled delivery, since more complex metal chelation environments can be generated throught the rational combination of the single decoration motifs kwnown to be effective for the capture of specifications.

Acknowledgments This work has been financially supported by “Ministerio de Economía y Competitividad” MAT2016-76739-R (AEI/FEDER, EU), and the “Gobierno Vasco” Basque University Research System Group, IT-630-13, which we gratefully acknowledge. The European Commission Research & Innovation H2020-MSCA-RISE-2017 (Ref.: 778412) INDESMOF project is also acknowledged. ELKARTEK-ACTIMAT, ELKARTEK-LION, ELKARTEK-LISOL, HAZITEK-SIMAN, and PIBA-LIMOFILM (PIBA-2018-06) projects funded by the Basque Government Education and Industry departments are also gratefully acknowledged.

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CHAPTER FOUR

Multifunctional materials for clean energy conversion R. Krishnapriya1, Devika Laishram1, M. Shidhin2, Rakesh K. Sharma1 1

Sustainable Materials and Catalysis Research Laboratory, Department of Chemistry, Indian Institute of Technology Jodhpur, Jodhpur, Rajasthan, India The Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Chennai, Tamil Nadu, India

2

4.1 Introduction With the rapid depletion of fossil fuels, rising environmental concerns, and population growth, it is inevitable to develop clean energy technologies to power our future society [1e4]. These energy conversion and storage technologies are anticipated to be sustainable and also capable of meeting our long-term energy needs. During the past few years, extensive research interests have been devoted to the advancement of energy conversion devices, as they play a crucial role in the prosperity and economic growth of a country. Particularly, the energy conversion technologies such as solar and fuel cells have proved to be highly reliable and can offer clean and sustainable energy at affordable rate [5e8]. However, the performance potential of these devices, such as output voltage, conversion efficiency, and stability, are greatly relied on the materials used. The energy conversion process comprises physical and/or chemical reactions at the surface of the material or at its interfaces; hence, the specific surface area, surface energy, and even surface chemistry can critically affect the device performance. Recently, the use of nanomaterials as the building blocks to fabricate energy conversion devices has opened up a novel path for the development of highly efficient and robust electrode materials with required properties [9,10]. These materials exhibit exceptional electrical conductivity, thermal conductivity, dielectric constant, and mechanical quality factor owing to the quantum dimension effects and macroscopic quantum tunneling effects. The lower dimensions of nanomaterials offer satisfactory mass, heat, and charge transfer through the advantage of accommodation of dimensional variations in accordance with chemical reactions and phase transitions. Among different nanomaterials, multifunctional nanomaterials are receiving Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00007-X

© 2021 Elsevier Ltd. All rights reserved.

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specific attention. These materials can show distinct optical, magnetic, and electrical properties, which are anticipated to offer enhanced performance when applied in a device [11,12]. Generally, the materials that are designed to perform multiple roles via a prudent combination of several functions are called multifunctional materials. The research on these materials involves understanding the structure-property relationship that is leading to the design and fabrication. However, the fabrication of multifunctional material with desirable functions is relatively challenging, from searching for novel materials, understanding its properties, and tuning the properties in order to apply in a technological device. In this chapter, a comprehensive summary of various multifunctional nanomaterials at a broader perspective to overcome the challenges of energy conversion devices is presented.

4.2 Energy conversion devices 4.2.1 Solar cells There are different energy conversion devices that effectively use the nanostructured materials to convert the energy into a useful form. Among these, solar cells are the central energy conversion devices that deliver clean electrical energy, where the renewable solar energy is directly converted into electrical energy without releasing any greenhouse gases. The main features of solar cells are the internal and external quantum efficiency and energy conversion efficiency. Solar cells are mainly classified into three types. The first-generation silicon solar cells made of Si wafers and mainly single crystal-based on p-n junctions, which are the most widely used solar cells and which have taken over the current photovoltaic (PV) device market and have reported maximum efficiency of nearly 25% [13]. The secondgeneration PV devices use thin-film technologies, which include cadmium telluride (CdTe), copper indium gallium diselenide (CIGS), and amorphous thin-film silicon. However, the efficiency is much below that of silicon solar cells. Third-generation solar cells include solution processing technology of semiconducting organic macromolecules, inorganic nanoparticles, or hybrids [14]. Copper zinc tin sulfide solar cell, dye-sensitized solar cells (DSSC), organic solar cells, perovskite solar cells, and quantum dot solar cells (QDSC) come under this classification. A schematic representation of the different generations of solar cells and their achieved efficiencies are given in Fig. 4.1. Among different types of solar cells such as DSSCs, hybrid organic solar cells and QDSCs are based on the advances in nanotechnology [15e18].

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glass

Perovskite SCs 19.7%

4th Generation

HIT SCs 24.7%

Substrate

1st Generation

ETL AI

3rd Generation

2nd Generation

GaAs SCs 28.9%

DSSCs 11.9%

TCO HTL

CIGS SCs 21.7%

Multijunction SCs 38.8%

Inorganic in organic nano particle polymer hybrid

Monocrystalline Si SCs Polycrystalline Si SCs 26.7% 21.9%

Figure 4.1 Schematic diagram showing the evolution of different generations of photovoltaic solar cells with respective device structures and the current photovoltaic efficiencies. Reproduced with permission from Elsevier A. Sahu, A. Garg, A. Dixit, A review on quantum dot sensitized solar cells: past, present and future towards carrier multiplication with a possibility for higher efficiency, Sol. Energy 203 (2020) 210e239.

The solar energy harvesting and its conversion is achieved by nanostructured semiconductor interface with a dye in case of DSSC, a conjugate polymer in organic solar cells, and semiconductor nanocrystals for QDSCs [19]. Improving the efficiency of these devices involve bringing the effective photo-induced charge separation, as well as the transport of charge carriers across the nanoassemblies, which is the main challenge. The advance of nanoassemblies to improve light harvesting efficiency, thermodynamic and kinetic measures for effective device design, and approaches for effective utilization of photo-induced charge separation in donor-acceptor molecules to fabricate nanostructure-based solar cells are the hot topic of research in recent years [2,20e22].

4.2.2 Fuel cells Fuel cells are energy conversion devices that can convert chemical energy from a fuel into electrical energy via a reaction with oxygen or other

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oxidizing agents. Recently, fuel cells are getting considerable consideration as an alternative energy source owing to their high efficiency without the release of harmful chemicals to the surroundings. The basic structure of a fuel cell contains an electrolyte that is in contact with a porous anode and a cathode. The fuel comes into the anode, and then an oxidant moves in through cathode, which is divided by a selectively conductive electrolyte. Further, the conduction occurs through the electrolyte in either anode to cathode or cathode to anode direction. The main conductive charge carriers 2  comprise Hþ, CO2 3 , O , OH , etc. [23,24]. The important commercial fuel cell technologies are divided into many groups in accordance with the operating conditions and which is given in Fig. 4.2. The most efficient and widely used fuel cells applied in cellular phones and laptops are direct methanol fuel cells and proton exchange membrane fuel cells because of their high power density and lower operating temperatures [25e28]. However, there are some challenges faced by fuel cell technology, such as fabricating suitable electrodes for the flexible electronics, replacing the high-cost noble metal electrocatalysts, and preventing electrode poisoning. The major challenge associated with this technology is the cost of fuel. However, significant advances in the development of materials with improved properties can make this technology more reliable. Current research is focused on the development of high-efficiency fuel cells of low cost with ecofriendly nanostructured electrode materials [29,30].

Increasing operating temperature

Fuel cells

Polymer electrolyte(PEM)

Direct hydrogen(PEM)

Direct methanol (DMFC)

Alkaline electrolyte(AFC)

Molten phosphoric acid electrolyte (PAFC)

Molten carbonate electrolyte (MCFC)

Solid oxide electrolyte (SOFC)

lons conducted through electrolyte Proton(H+) Hydroxide ion(OH–) Oxide ion(O2–) Carbonate ion(CO32–)

Figure 4.2 Types of different commercially available fuel cells.

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4.3 Multifunctional nanomaterials for energy conversion applications The primary requirement of nanostructured materials for PV devices such as DSSCs and QDSCs is to have a large specific surface area for the anode to adsorb sufficient dye molecules and quantum dots (QDs). This can act as an antenna for light absorption. To achieve the maximum absorption of incident light, the narrow thickness of photoanode with minimum interface charge recombination is desired [31e33]. Moreover, the material must possess remarkable charge mobility and long lifetime, with some light scattering effect or photon trapping ability. Fuel cell technology faces significant challenges like low energy density, power density, and operation reliability along with the high cost. Nanostructured materials can overcome most of these challenges by providing high specific surface area and improved conductivity with low polarization. These materials have the potential to offer excellent nanoporous structures with multifunctional chemical properties for highly intrinsic electroactivity and excellent utilization [34]. Nanostructured semiconductor metal oxides of TiO2, ZnO, SnO2, Fe2O3, WO3, and BiVO4 were extensively explored as photoelectrode materials for solar cells owing to their tunable semiconducting properties, high stability, and easy fabrication. Mesoscopic TiO2 film is the main component of DSSCs, organic photovoltaics, and QDSCs, which helps to capture electrons from the excited sensitizer or quantum dot and allows the easy transport of electrons to the transparent collecting electrode [35,36]. The charge separation and transports occur at the photoelectrode/electrolyte interface or within the photoelectrode which is critical for device performance. These reactions mainly depend on the morphology and structure of metal oxide used. Thus, metal oxides of different morphologies have been applied to tailor optical and electronic properties. A systematic study on the influence of hierarchically structured morphology of different metal oxide (MO) photoanodes on the performance of DSSCs was reported by various research groups [37e39]. Sun and coworkers reported a novel multifunctional photoanode consisting of one-dimensional (1D) TiO2 nanotube arrays modified by TiO2 nanoparticles as the bottom layer and three-dimensional (3D) TiO2 submicron spheres as the top layer by a facile one-pot chemical bath deposition method. The prepared 1D-3D bilayer photoanode demonstrates a power conversion efficiency (PCE) of 6.93%, leading to a 26.5% increment of PCE compared to 5.48% by TiO2 bare nanotube arrays [40]. A high PCE of

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(a) Au

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FTO/glass

(b)

(c)

(d)

(e)

(f)

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Figure 4.3 (a) Schematics of the perovskite solar cell architecture with half of its interface covered by cp-TiO2 ETL and the other half covered by TiO2 nanocolumn arrays.

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9.3% was reported by single crystallike anatase TiO2 nanowires network prepared by an “oriented attachment” mechanism. The resulting nanowire (NW) shows the high rate of electron transfer through the TiO2 film [41]. The effect of length of nanotube can also influence the PCE of DSSC which was reported by Liu et al. Nanotubes (NTs) with length up to 20.8 mm were achieved which improved the PCE [42]. Recently, vertically aligned 1D TiO2 nanocolumn arrays were fabricated by glancing angle deposition to apply as electron transport layer (ETL) to fabricate functional triple-cation lead perovskite halide solar cells based on Cs0.05(FA0.83MA0.17)0.95Pb(I0.83Br0.17)3 [43]. The schematic representation of the device architecture and the scanning electron microscopy (SEM) images of TiO2 nanocolumn are given in Fig. 4.3. A promising PCE of 16.38% with 5% boost of short-circuit current density and a 7% improvement in its power conversion efficiency, together with a significantly prolonged shelf life, were obtained by careful tuning of the TiO2 column properties. Another potential MO for DSSCs are ZnO and SnO2; they have a energy band position and physical properties comparable to TiO2 but these materials have higher electron mobility, which is capable of improving the electron transport efficiency leading to reduction in recombination loss. Morphology engineering of ZnO, doping, heterostructures of ZnO/TiO2, phase and dimensionally controlled synthesis of oxides at subzero temperatures, etc., succeeded in improving the efficiency of solar cells [44e49]. The ZnO morphology-dependent PCEs of DSSCs with respect to various reaction conditions are given in Fig. 4.4. ZnO demonstrated as a promising ETL to TiO2 for perovskite solar cells (PSCs) exhibiting remarkable efficiency of 18.9% [50]. Recently, ZnO with a high crystallization multiple cation perovskite absorber showed a high efficiency of over 20% [51]. Compared to TiO2 photoanodes, SnO2-based DSSCs are showing less PCE. Zhang et al. fabricated SnO2@Air@TiO2 hierarchical urchinlike double-hollow nanospheres as

(b, c) Top-view SEM images of the cp-TiO2 layer (b) and the TiO2 nanocolumn arrays (c). (d, e) SEM image of the perovskite film deposited on a cp-TiO2 ETL (d) and on TiO2 nanocolumn arrays (e). (f, g) Cross-sectional SEM image after the deposition of a hole transport polymer on the part of the sample with only a cp-TiO2 ETL (without nanocolumns, (f) and the part of the sample with TiO2 nanocolumn arrays (g). Reproduced with permission from American Chemical Society Z. Hu, J.M. García-Martín, Y. Li, L. Billot, B. Sun,
lez, L. Aigouy, Z. Chen, TiO2 nanocolumn arrays for F. Fresno, A. García-Martín, M.U. Gonza more efficient and stable perovskite solar cells, ACS Appl. Mater. Interfaces 12 (5) (2020) 5979e5989.

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Time dependent

Power dependent

Temperature dependent

(b)

(f)

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(k)

8 7 6

125°C 50°C

(d)

Ambienf

PCE = 1.96%

5

300 W 200 W 50 W

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(h)

PCE = 0.26% PCE = 0.15%

1 0.1

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0 0.0

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(I) PCE = 2.58%

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Photocurrent density (mA/cm2)

(i)

Photocurrent density (mA/cm2)

(e)

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(a)

0 0.0

PCE = 1.18% PCE = 0.69%

0.1

0.2

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Voltage (V)

Figure 4.4 A comparative study of morphology-dependent J-V curves of hierarchical ZnO flower growth stages illustrated in FE-SEM images. Reproduced with permission from Royal Society of Chemistry R. Krishnapriya, S. Praneetha, A.V. Murugan, Investigation of the effect of reaction parameters on the microwave-assisted hydrothermal synthesis of hierarchical jasmine-flower-like ZnO nanostructures for dye-sensitized solar cells, New J. Chem. 40 (6) (2016) 5080e5089.

photoanode. This composite structure of DSSC demonstrated a maximum efficiency of 6.77% owing to the increased transmission of photogenerated electrons [52]. Composite and hybrid photoanodes are another promising alternative for improved performance of DSSCs. GO/SnO2 hybrid nanocomposite-based photoanode for DSSCs have achieved a PCE of 8.3% with higher dye-loading, rapid electron transport, superior light scattering, and lower electron recombination rate [53]. Doping of lithium ions is proved to be the best way to improve the PCE of SnO2-based solar cells.

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Park et al. reported the fabrication of a highly stable and efficient electron transporting layer using Li-doped SnO2 (Li:SnO2) relatively at low temperature in wet-chemical route [54]. The doped Li in SnO2 improved conductivity and favored a downward shift of the conduction band minimum of the oxide, which enabled easy path for the electrons injection and transfer from conduction band of perovskite. An exceptionably high PCE of 18.2% and 14.78% were achieved for the rigid and flexible substrates, respectively, which is shown in Fig. 4.5. Perovskites are naturally occurring minerals of CaTiO3, and the materials which have same type of crystal structure as CaTiO3 are generally called perovskite materials [55]. These materials possess a cubic or tetragonal crystal structure with the stoichiometric formula AMX3, where both A and M are metal cations with A larger than B and X is the anion, usually oxides or halogens [56]. Commonly, each cation M is octahedrally coordinated with the anion X to form MX6 octahedra (Oh), which is the basic building block of the perovskite structure. These Oh units are connected in a 3D corner-sharing configuration, with the cation A surrounded in the space formed between adjacent MX6 octahedron to neutralize the charge of the structure. The unique physical properties of perovskite materials such as high-absorption coefficient, long-range ambipolar charge transport, low exciton binding energy, high dielectric constant, ferroelectric properties, etc., make this material a suitable candidate for photovoltaic applications. Distortions exist in the perovskite structure that results in oxygen nonstoichiometry comprising both oxygen deficiency and oxygen excess. Thus, perovskite oxides offer great flexibility in redox active sites, oxygen vacancies, and physicochemical properties with tunable compositions due to the numerous possible substitutions at both A and B sites. Perovskite oxides can exist in different structures such as double perovskites and layered perovskite. Owing to its varied compositions and structure, the material offers excellent thermal stability, redox properties, oxygen mobility, and electronic and ionic conductivity; perovskites became very attractive in solar cells around the world. With the advantage of unprecedentedly superb optoelectronic properties, together with a carrier diffusion length of up to w100 mm and a low exciton binding energy of w20 meV, this material enables the formation of free carriers even at room temperature, high optical absorption coefficient of w105 cm1, and low nonradiative recombination for optoelectronic device applications. PSCs have shown more impressive progress than other PV technology with the certified efficiency of 25.2%. However, the poor stability of the perovskite

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(b)eV

(a) Au

3.95

4.22

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4.35

4.5

Perovskite

CH3NH3PbI3

Li:SnO2

SnO2

5.4

FTO

Li:SnO3

8.17

FTO

8.3

300 nm

(d) 25

20

20

15

Current density (mA/cm2)

25

Backward (15.29%) Forward (13.55%)

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EQE (%)

80 60 40

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Li:SnO2

20 0 300 400 500 600 700 800 900 Wavelength (nm)

(f) Number of devices

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PCE (%)

Figure 4.5 (a) Cross-sectional SEM image of a PSC. (b) Energy diagram of FTO, ETLs, and perovskite (CH3NH3PbI3). (c, d) J-V curves measured at backward-forward scan for 200 ms of scan-delay time. The measurements were performed under simulated AM 1.5 G sunlight of 100 mW/cm2. (e) EQE spectra of the devices comprising SnO2 and Li:SnO2 as an ETL, respectively. (f) Histogram of PCEs for 30 devices. Reproduced with permission from M. Park, J.Y. Kim, H.J. Son, C.H. Lee, S.S. Jang, M.J. Ko, Low-temperature solution-processed Li-doped SnO2 as an effective electron transporting layer for highperformance flexible and wearable perovskite solar cells, Nano Energy 26 (2016) 208e215.

materials with regard to humidity, heat, light, and oxygen is the main challenge that should be addressed before its commercialization. Recently, CsYbI3 cubic NCs presented strong excitation-independent emission and high photoluminescence quantum yields of 58% [57].

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[HC(NH2)2]0.83Cs0.17Pb(I0.6Br0.4)3, Csþ-Doped 2D (BA)2(MA)3Pb4I13, halide perovskites, with a typical structure of ABX3 (A ¼ CH3NHþ 3, þ 2þ 2þ  CH(NH2)þ , Cs ; Btheir achieved efficiencies ¼ Pb , Sn ; X ¼ Cl , 2   Br , I ), and a-CsPbI3-based PSCs showed promising PCE [58,59]. Additionally, perovskite oxides with high ionic/electronic conductivity find good application as anode, cathode, and electrolyte in solid oxide fuel cells (SOFCs) [60,61]. This material is the practical choice at operation at 700e900 C because of high electrochemical activity for the O2 reduction reaction, high thermal stability, and compatibility. It shows excellent microstructural stability and long-term performance stability. Perovskite oxides with nanoporous nature exhibited greater powder density and lower area-specific polarization for fuel cell applications [62]. Novel nanostructured carbon-based metal-free electrode materials can be used as superior catalyst supports for energy conversion devices due to their remarkable electric conductivity. These materials can act as efficient photo-/electrocatalysts to enable the crucial chemical reactions. These specific reactions include the oxygen reduction reaction (ORR) in fuel cells and the iodine reduction reaction in DSSCs. The incorporation of pblock elements such as N, P, Si, and B, as well as the heteroatoms (O, S, Se, F, Cl, Br) can greatly improve the device performance. A 3D holey N-doped graphene of the special hierarchical framework structure, high specific surface, and rich N species is used for ORR and OER catalysis. The performance was found to be comparable to the commercial RuO2 and Pt [63]. The peculiarity of carbon structures is the intrinsic defects existing in carbon structures, which can induce charge transfer and density of state change, thus creating more active centers for reactions. Various nanoarchitectures, with a range of sizes, shapes, compositions, and structures, have shown good potential to catalyze the reactions in solar cells and fuel cells. Several carbon-based nanostructures, such as fullerenes, nanotubes, and graphene are used in solar cells [64,65]. Graphene is most commonly used for fabrication in carbon-based organic photovoltaic cells due to its excellent electron transport properties and extremely high carrier-mobility. In addition, its energy level can be tuned easily through controlling its size, layers, and functionalization. These materials can be used for the fabrication of transparent conducting electrodes, as composite for semiconducting layer,

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in electrolyte and as counter-electrodes. These materials are established as efficient candidates to replace or modify the existing components in solar cells [66]. Carbon nanotubeebased PSCs can exhibit higher high-temperature operational stability than those of metal-based electrodes. Interesting report on ex-situ vapor-assisted doping of CNT electrode for efficient and stable PSCs is shown in Fig. 4.6. The enhanced conductivity with a favorable energy alignment of modified CNT is able to achieve a PCE of 17.56% with greater stability. As the method employed was facile, rapid, and readily applicable to large area, module process can successfully offer a solution to overcome the efficiency limit in CNT-laminated PSCs [67]. In the photoanode, graphene materials have resulted in improved photocurrent. Additionally, an atomically dispersed Co-doped carbon catalyst with a core-shell structure was developed via a surfactant-assisted metale organic framework approach and applied to fuel cells. The CoN4 active sites in catalysts revealed unprecedented catalytic activity for ORR with a halfwave potential of 0.84 V and promising stability in challenging acidic media [68]. N-Doped carbon nanobubbles, spheres, and their graphene composites have also showed long-term operational stabilities and comparable gravimetric power densities when applied in various energy device applications [69,70]. Owing to the advantages like low cost, earth abundance, and ease of synthesis and practical applications, metal chalcogenides (MC) have found great importance for application in energy conversion devices. These nanostructured materials have found to be an effective alternative to Pt/C electrocatalysts in fuel cells. Baresel et al. reported the electrocatalytic activity of thiospinels and other sulfides for oxygen reduction in acidic electrolytes [71]. Later, chevrel-phase Mo4Ru2Se8 chalcogenides and related compounds were reported and showed interesting catalysis properties [72]. Various MC of cobalt, nickel, and iron show unique electrical properties, excellent catalytic activities, and outstanding stabilities. These materials were successively used as counter-electrodes (CE) for DSSCs. Electrochemically deposited CoS NPs on flexible conducting oxide films have reported 6.5% of PCE [73]. Accordingly, sulfides and selenides of iron, cobalt, and nickel (FeS2, FeSe, Co0.85Se, and NiSe-Ni3Se2) have been reported extensively as potential low-cost CE materials [74e76]. Chang et al. prepared

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Figure 4.6 (a) Schematic illustration showing structure of planar heterojunction perovskite solar cells based on TFMS doped CNT (b) Current density-voltage (J-V) curves of the highest efficiency perovskite solar cells incorporating bare CNT and CNT doped by TFMS (30 and 50 s of doping time). (c) EQE spectra and corresponding integrated short-circuit current density (JSC) of the device based on bare CNT and CNT doped by TFMS for 30 s (d) JSC, (e) open circuit voltage (VOC), (f) fill factor (FF) and (g) power conversion efficiency (PCE) of the perovskite solar cells incorporating bare CNT and CNT doped by TFMS. Reproduced with permission from American Chemical Society J.W. Lee, I. Jeon, H.S. Lin, S. Seo, T.H. Han, A. Anisimov, E.I. Kauppinen, Y. Matsuo, S. Maruyama, Y. Yang, Vapor-assisted ex-situ doping of carbon nanotube toward efficient and stable perovskite solar cells, Nano Lett. 19 (4) (2019) 2223e2230.

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Co9S8 nanocrystal-based nanoinks to fabricate uniform, crack-free Co9S8 thin films with a spray deposition technique, and when applied as counter electrode (CE) for DSSC showed PCE of 7.02  0.18% under AM 1.5 solar illumination [77]. The morphological features and the photoelectrical and electrochemical characteristics of the synthesized material are shown in Fig. 4.7.

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Figure 4.7 (a) Schematic diagram of Co9S8 nanocrystals synthesized on FTO by spray deposition; (b) Photograph of Co9S8 nanoink used for the thin-film deposition; (c) Photograph of the Co9S8 nanocrystal thin film deposited on a 10 cm  10 cm FTO substrate; (d) SEM image of the cross-sectional and plan (inset) views of the Co9S8 film/FTO; (e) Photograph of a simulated DSSC cell with 2 cm2 working area; (f) Photocurrent density-voltage plots of DSSCs based on platinum, Co9S8 on FTO, and Co9S8 on Mo CEs obtained under AM1.5 illumination; (g) Nyquist curves of the DSSCs based on platinum, Co9S8 on FTO, and Co9S8 on Mo CEs. Reproduced with permission from American Chemical Society.

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Moreover, along with monometallic chalcogenides, binary and ternary MC have been successfully applied as CE materials for DSSCs because of their exceptional electrochemical properties. By the careful control over chemical composition, ternary and polyphyletic compounds based on iron, cobalt, and nickel, such as NiCoSe, Co-Fe-Se/S, and NiCo2O4, provided catalytically active CE with good electronic conductivity [78e80]. Bimetal transition metal alloys and compounds such as MIn2S4 (M ¼ Fe, Co, Ni), NiCo2S4, CoFeS2, and (Ni, Fe) S2 also successfully reported and exhibited high performances due to the coexistence of two different cations in a single crystal structure [81,82]. Among them, siegenite (NiCo2S4) with a normal thiospinel crystal structure unveils outstanding electrocatalytic performance due to the synergistic interactions between Co2þ/3þ and Ni2þ. The hierarchical nanostructured NiCo2S4 with urchinlike, porous sheetlike, hexagonal, and flowerlike morphologies has been extensively reported for Pt-free DSSCs. However, the use of pristine materials as CEs is facing metal aggregation problems; so its application in DSSCs is limited. A possible solution to this problem is to make a composite with graphene. Graphene is particularly interesting because of its layers of two-dimensional (2D) sp2-bonded carbon sheets, its unique structure, and exceptional physical properties, such as high electrical and thermal conductivities, mechanical flexibility, charge transport mobility, huge specific surface area, good chemical stability, and optical transparency. Thus, nanocomposites of carbon nanotube and graphene-integrated transition metal-based electrocatalysts have abundant research consideration. Accordingly, the growth of nanocrystalline NiCo2S4 on graphene nanosheets (GNSs) has been reported as the GNSs serve as 2D electrical conductive scaffold to electroactive NiCo2S4 and also are capable of enhancing charge transport factors, thereby refining the electrochemical performance. A sustainable rapid microwave-solvothermal (MW-ST) synthesis approach to develop NiCo2S4 nanocrystals and their nanohybrids with GNS was reported with a PCE of 7.98% [83]. The morphological features and the PV performance of the CE are given in Fig. 4.8.

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Figure 4.8 FE-SEM images of NiCo2S4 nanocrystals with (a) nanocrystalline aggregates (NCS-1), (b) the NiCo2S4egraphene hybrid (NCS-1/GNS), (c) tremella-like NCS-2, and (d) porous bead-cum needlelike NCS-3 and prepared by the MW-HT/ST method. The inset figures show the magnified FE-SEM images of the respective samples. (e) A comparative plot of JeV curves of DSSCs fabricated with different morphologies of pristine NiCo2S4, the NCS-1/GNS hybrid, Pt and pristine GNS counter-electrodes and a comparative histogram of power conversion efficiencies (PCEs) versus respective counter-electrodes (CEs). (f) Comparative IPCE spectra of DSSCs fabricated with different morphologies of pristine NiCo2S4, NCS-1/GNS hybrid, Pt and pristine GNS counter-electrodes showing high photoresponse. Reproduced with permission from Royal Society of Chemistry R. Krishnapriya, S. Praneetha, A.M. Rabel, A. Vadivel Murugan, Energy efficient, one-step microwave-solvothermal synthesis of a highly electrocatalytic thiospinel NiCo2S4/graphene nanohybrid as a novel sustainable counter electrode material for Pt-free dye-sensitized solar cells, J. Mater. Chem. C 5 (12) (2017) 3146e3155.

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4.4 Conclusion Energy conversion devices require efficient materials which possess distinctive electrical, optical, mechanical, and thermal properties. Recently, the development of nanoscience and technology has helped to meet the requirement of many materials for full application impact. During the past few years many multifunctional nanomaterials have been developed for energy applications. The advancement in this area is achieved by the properties such as excellent electrical and thermal conductivity, exceptionally large surface area, and chemical stability. The main challenges behind the multifunctional material are the higher performance limit, fewer functions, toxicity, availability, and high cost. Deep understanding of structure-property-relationship with the help of advanced characterization techniques is required in order to realize the large-scale commercial application of these materials. Although the energy conversion device performance using these materials enhanced considerably, to meet the future energy demands more developments are needed.

Acknowledgments Authors acknowledge financial support from PAN-IIT DBT center for Bioenergy grant number BT/EB/PANIIT/2012, Indian Institute of Technology Jodhpur for facilities and infrastructure support. We also acknowledge Dr. Kiran P Shejale for his contribution on research on DSSC.

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CHAPTER FIVE

Advances in thermochromic and thermoelectric materials Ahmed Esmail Shalan1, 2 Nikola Pe
rinka1, Esraa Samy Abu Serea1, Mohamed Fathi Sanad3 1

BCMaterials-Basque Center for Materials, Applications and Nanostructures, Martina Casiano, UPV/EHU Science Park, Leioa, Spain 2 Central Metallurgical Research and Development Institute (CMRDI), Helwan, Cairo, Egypt 3 FabLab, Centre for Emerging Learning Technologies (CELT), and Electrical Engineering Department, The British University in Egypt (BUE), Cairo, Egypt

5.1 Introduction “Smart materials” are materials that can modify their configuration, properties, as well as functions in response to different surrounding parameters, such as temperature, pressure, and magnetic and electric fields, often used as actuators and sensors [1e4]. Mostly, “smart” or “intelligent” materials are embedded into systems whose fundamental features are being auspiciously transformed to provide a desired performance [5]. Smart materials present certain advantages and disadvantages. On the one hand, the benefits consist of high energy density, fast response, and fewer moving parts. On the other hand, the shortcomings consist of restricted sensitivity to harsh environmental situations [6,7]. In addition, one of the main advantages of smart materials regarding different applications is their nonlinear properties via hysteresis effects [8,9]. Contemporary progress in smart materials appeared when these materials were incorporated in conventional devices, e.g., actuators and sensors, and displayed the basic active properties of smart or intelligent structures [10]. Furthermore, external disturbances and additional responses of smart structures toward several parameters, such as temperature, pressure, and strain, with active control in real time, provided actions to identify problems and meet anticipated needs [11,12]. In the case of an adaptive structure, actuators can change system status, whereas sensors can allow continuous monitoring of the state of the structure. Consequently, by conjoining both sensory and adaptive structures, we can obtain controlled structures. Thus controlled structures can be Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00013-5

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produced by combinations of sensors and actuators, which can accomplish both structural and control functions, together with control logic and electronics [13e17]. Several possible applications of smart materials allow the controlling of vibration and shape features, such as micropositioning, antenna shape control, and automatic flow control valves. Shape control and active suspension systems for vehicles and active vibration control systems in aircraft use smart materials as their desired resources [18e21]. Other applications involve taking advantage of thermal variations, such as thermochromic and thermoelectric materials.

5.1.1 Thermochromic smart materials Thermal energy storage classifications are considered the way to maintain energy and decrease the environmental effect of energy usage [22e28]. To gain further benefits from thermal energy storage, the practice of latent heat storage is considered an effective method, which has the characteristics of high heat storage density and small differences in released heat [29]. Chromic colorants can change, radiate, or erase color because of their orientation by exterior stimuli. “Chromic” as a suffix denotes reversible change of color by extension of other physical possessions. Outside stimuli can be light (photochromic), heat (thermochromic), electricity (electrochromic), pressure (piezochromic), liquid (solvatochromic), or electron beam (carsolchromic) [30e34]. The usage of chromic colorants has attracted attention from science, technology, materials, and fashion, and has allowed the development of novel research for numerous applications [35e37]. There is a substantial need for the application of thermochromism in the field of smart materials, which are engineered to sense and respond to external environmental surroundings [38]. As a result of the opportunities to improve novel innovative designs, color-changing smart materials are producing awareness among artists and designers because of their interaction, responsiveness, and ultimate functionality. Fig. 5.1 is a schematic diagram showing the different components of a thermochromic device module structure. Accordingly, color-changing tools suggest exceptional and inspiring design opportunities to the designer. In the last few decades, various scientific works have been available regarding chemistry and the mechanisms of thermochromic colorants that provide important data about the preparation, characterization, and possible application of these materials [39,40].

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Printed thermochromic layer Printed resistive heating elements

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Figure 5.1 Schematic diagram showing the different components of the thermochromic device module structure.

5.1.2 Thermoelectric smart materials A thermoelectric device consists of two different types of semiconducting materials, p-type containing holes and n-type containing electrons, which are connected together with a strong junction to facilitate the carrier movement [41]. In addition, the temperature difference at the junction is the main reason behind the flow of carriers away from the junction, leading to an open-circuit voltage and accordingly creating an electrical generator. This phenomenon is termed the Seebeck effect [42]. On the other hand, the Peltier effect is designed when an electric current is transient over the junction, while both electrons and holes move away from the junction and carry heat energy away to cool the junction down. Furthermore, Fig. 5.2 shows a schematic diagram of the different components and arrangement of a typical wearable thermoelectric module [43]. Through the transport system, the Seebeck coefficient (S), electrical conductivity (s), and electronic thermal conductivity of a material can be

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considered depending on the Boltzmann transport equation. This is considered as the physical derivation of each transport coefficient, and the effects of temperature and carrier concentration, to optimize the power factor (S2s), and discover probable directions for additional improvement to the power factor [44]. Thermoelectric materials are able to transmute a temperature gradient into electrical energy when a temperature difference occurs between the hot and cold terminations of the substrate [45]. Thermoelectric devices are used because they are easily scalable for energy conversion and have no moving parts, which enable their application from small-scaled synthesis to huge industrial facilities [46e48]. With the help of nanotechnology, nanostructures are applied to maintain acceptable electrical conductivity as well as higher voltage output thermoelectric devices. The application of nanostructured materials and low-dimensional superlattice morphologies has led to enhancement of the figure of merit (zT). As an example, thermoelectric materials based on bismuth telluride (Bi2Te3), zinc antimonite (ZnSb), and silicon germanium (SiGe) are compounds with the best zT, and are excellent models that emphasize the function of these materials for thermoelectric systems [49]. An ideal thermoelectric material will thus show low thermal conductivity but high electrical conductivity [50]. In this chapter, we will focus on the different materials to fabricate effective thermochromic as well as thermoelectric devices and focus on the appropriate pathways for the desired applications of these techniques.

5.2 Thermochromic and thermoelectric materials and devices Thermal losses are considered the foremost source of the significance of harvesting outstanding energy, which depends mainly on the existence of those losses. Typically, such energy losses are linked to industrial progressions and unconstrained to an environment that is deprived of any kind of possible recovery. There are different types of materials that depend on these kinds of energy losses, and thermochromic as well as thermoelectric materials are considered probable examples of those materials. Thermochromic materials can transform their color in response to a change in temperature, and are considered as either reversible or irreversible. Reversible thermochromes may be used in temperature-sensing applications that indicate the current temperature, while irreversible thermochromes may be used to provide a temperature limit [51e53]. There are several

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technological advantages for using thermochromic materials. It is also considered that sensor materials are a kind of smart material, and may be used with other technologies [54e56]. Even though different researchers have discussed the phenomenon of thermochromism previously [57e60], most of the applications are limited to the dissolved state of organic compounds and metal complexes compared to the solid state of these compounds [61]. Additionally, the color of these compounds is obtained from structural colors like photonic liquid crystals [62]. Other isolated examples of thermochromism, depending on isomerization reactions in the solid state, have also been studied [63e65]. On the other hand, for thermoelectric materials, different efforts by many researchers have attained high-performance nanostructured materials to enhance the effectiveness of conventional thermoelectric converters as well as obtain a new direction for those materials to generate electricity through the capture of residual energy in household appliances [66e70]. Furthermore, different studies have shown improvements in the progress of thermoelectric materials with rare earth elements, reaching efficiencies higher than 20% [71e75]. Additionally, products such as drinking fountains, cellars, incubators, organ transport boxes, vaccine storage chambers, and portable pharmaceutical refrigerators, which use electricity for refrigeration and heating of containers, are commercially available for the Peltier effect [76e78]. The use of thermoelectric modules is limited to small volumes and localized cooling to meet their technical and economic feasibilities [79e81].

5.3 Applications of thermochromic materials and devices More attention has been given to thermochromic materials because of their capability for color variation as a function of temperature, discovered as a result of the steady color change above a wide temperature range because of the deviations in the widths of the absorption bands [81]. The interaction phenomena of the transition metal with a suitable donor solvent are attributable to the energy alternation of the dd transitions and configuration changes and can be explained by the behavior of thermochromic actions that appear via numerous transition-metal complexes, for instance, nickel, copper, and halide complexes [82e84]. In addition, reversible color change, which occurs in response to temperature distinction via heating or cooling, can characterize or explain the phenomenon of thermochromic behavior

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[85]. Additionally, there are several applications that depend on thermochromic systems, such as energy [86], flexible electronics [87], the textile industry [88], and smart food packaging [89], which can be applied in different capacities.

5.3.1 Applications of thermochromic polymers in protected quick response codes A thermochromic fluorescent as a rapid and reversible supramolecular polymer hydrogel was prepared through orthogonal self-assembly, including hosteguest connections and contacts with metal coordination. This supramolecular polymer hydrogel displays awareness toward various incentives that may encouraged thru reasonable ligands, temperature and outstanding self-healing features afterward being damaged via external forces. In addition, after preparation of the fluorescent supramolecular polymer hydrogel, a three-dimensional quick response (QR) code was then constructed using this blue light-emitting, which could hide information under natural light at high temperature to enhance durability and security for a more protected state, or display it under UV light irradiation at low temperature in the presence of competitive ligands. Moreover, the protected QR code indicates exceptional self-healing performance even after being destroyed because of the dynamic reversibility of noncovalent interactions in the attained supramolecular polymer hydrogel [90]. Hydrogels are delicate, water-swollen, 3D polymer arrangements that can be developed by covalent or noncovalent bonds among the polymer chains [91e93]. Hydrogels have attracted much consideration because of their wide application in contact focal points, platforms for tissue building, medication conveyance and false ligaments, agribusiness, water remediation, and supercapacitors [94]. Notwithstanding, normal hydrogels with covalent bonds experience the effects of poor mechanical features attributed to their physical construction and high water content, restricting the extent of their applications. Most recently, different sorts of supramolecular noncovalent connections, for example, electrostatic, hydrophobic, hydrogen bond, metaleligand coordination, biomimetic, and stereo complexation [95e97], have been used in the development of hydrogels. Due to their superb boost responsiveness (toward ions [98], light [99], temperature [100], pH [101], and so forth), self-mending properties, particular biocompatibility and biodegradability, high rigidity, and high break toughness [102], supramolecular hydrogels have been generally applied in tissue engineering [103], injectable medication carriers [104], shape memory materials [105], and other fields.

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Current techniques to integrate supramolecular hydrogels mainly depend on solitary noncovalent communication, prompting materials that show just a solitary property or require complex blend steps, constraining their application. Consequently, there is a requirement for the improvement of simple strategies for building supramolecular hydrogels with different capacities. Various leveled, symmetrical self-assembly systems dependent on numerous noncovalent connections do not just enhance the types and elements of supramolecular hydrogels, but lead to progressively responsive supramolecular hydrogels [106]. The introduction of tetraphenylethylene (TPE) and its derivatives groups, also called aggregation-induced emission materials, inside the supramolecular hydrogel can successfully elude fluorescence quenching initiated through gelation, compared with traditional fluorescent dyes; they can also provide materials with bright blue fluorescence via their aggregated state [107,108]. Furthermore, a unique approach has been reported to create a fluorescent supramolecular hydrogel based on the hierarchical orthogonal self-assembly of TPE-bridged bis-(b-cyclodextrin) (b-CD)/adamantane (Ad), including hosteguest interactions and terpyridine-based metal coordination interactions (Fig. 5.3(a)). The resultant supramolecular hydrogel shows responsiveness to different boosts, instigated by various ligands and temperatures, as well as astounding self-mending properties subsequent to being damaged by external powers. Correspondingly, it radiates blue fluorescence, and the fluorescent power can be directed by temperature, allowing the hydrogel to be utilized as a light-emitting material with self-recuperating properties. At that point, a certified QR code was created (Fig. 5.3(b)). The QR code data are in a certified state and must be perused under explicit conditions. Moreover, the damaged QR code could be recovered because of the dynamic idea of noncovalent communications in the supramolecular hydrogel, and the repaired QR code could show the first data again under UV light [90]. In addition, the supramolecular hydrogel prepared via the solegel process shows strong emissions because of the aggregation of the TPE groups, which are monitored through the deviations in fluorescence intensity. It returned to its original state via cooling or through the consecutive addition of Zn2þ (Fig. 5.4(a) and (b)) [90]. The supramolecular hydrogel of b-CD/ Ad hosteguest interactions and terpyridine/Zn2þ metal coordination

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Figure 5.3 (a) A schematic representation of the structures of the host molecule H and guest molecule G, and the formation of a supramolecular hydrogel. (b) A cartoon representation of self-healing and information display or concealment of a supramolecular hydrogel QR code. Reproduced with permission from B. Li, C. Lin, C. Lu, J. Zhang, T. He, H. Qiu and S. Yin, Mater. Chem. Front. 4 (2020) 869e874. Copyright © 2020 Royal Society of Chemistry.

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Figure 5.4 (a) Cyclic switching of the fluorescence intensity of the H_G_Zn2þ hydrogel upon repeatedly adding cyclen and Zn2þ. (b) Temperature dependence of H_G_Zn2þ hydrogel spectra. (c) Photographs of the reversible solegel transition of the H_G_Zn2þ hydrogel. (d) Temperature dependence of dynamic shear moduli (s ¼ 1.0 Pa, o ¼ 10.0 rad/s) for the H_G_Zn2þ hydrogel. (e) Temperature dependence of the viscosity of the H_G_Zn2þ hydrogel. (f) Storage modulus (G0 ) and loss modulus (Gʺ) values of the H_G_Zn2þ hydrogel during continuous step strain measurements. The scanning frequency (o) was fixed at 10.0 rad/s. The H_G_Zn2þ hydrogel was subjected to 100% strain for 60 s, then back to 0.1% strain for 300 s; four cycles were performed. Reproduced with permission from B. Li, C. Lin, C. Lu, J. Zhang, T. He, H. Qiu and S. Yin, Mater. Chem. Front. 4 (2020) 869e874. Copyright © 2020 Royal Society of Chemistry.

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interactions displayed awareness to multiple stimuli due to their dynamic reversibility. Likewise, the hydrogel was disassembled by increasing the temperature, and was then transferred to the initial state when it was cooled to room temperature (Fig. 5.4(c)). Furthermore, the thermoreversibility of the supramolecular hydrogel had additional consideration as the dynamic shear moduli were estimated via utilizing the prepared arrangements inside a temperature range between 20 and 40 C [90]. In addition, at the point when the temperature was below 20 C, G0 was constantly bigger than Gʺ, affirming the formation of the gel state. However, at the point when the temperature was expanded to more than 22 C, both G0 and Gʺ diminished significantly, and Gʺ diminished quicker than G0 , which was ascribed to the solegel stage process. The state change of the supramolecular hydrogel was identified at around 25 C when the G0 /Gʺ hybrid point was noticed (Fig. 5.4(d)). Moreover, the solegel procedure of the hydrogel was achieved by thickness temperature oscillatory rheology utilizing a watery arrangement of the supramolecular hydrogel [109]. As shown in Fig. 5.4(e), when the temperature increased from 20 to 40 C, the consistency of the supramolecular hydrogel diminished significantly, showing the state change of the supramolecular hydrogel. These phenomena additionally affirmed the thermoresponsiveness of the supramolecular polymer hydrogel. As a target of that study to compare the traditional bar codes with the QR code, it was found that the QR is a two-dimensional durable code in a square shape, which can secure and store additional data and characterize extra information, which is convenient [110]. Because of the dynamic reversibility of their noncovalent cooperation, supramolecular hydrogels with remarkable responsiveness to different improvements and selfrecuperating properties are promising materials for the planning of QR codes. Such hydrogels could improve toughness and security of certified QR codes. In view of the thermally actuated solegel change property of supramolecular hydrogels, a three-dimensional certified QR code was effectively arranged in a defined form. As shown in Fig. 5.5(a), the QR code data were simply shown under the illumination of a UV light, and the data were hidden by an expansion in temperature. As a result of execution of the supramolecular hydrogel, the certified QR code likewise showed brilliant self-mending execution subsequent to being decimated. As shown in Fig. 5.5(b) and (c), the figure from the QR code vanished completely after 3 min, and the data were read out effectively and could be shown under UV light illumination.

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Figure 5.5 (a) Photograph and fluorescence images showing the display and concealment of QR code information. Supramolecular hydrogel QR code after damage: (b) at 0 min; and (c) at 3 min. Reproduced with permission from B. Li, C. Lin, C. Lu, J. Zhang, T. He, H. Qiu and S. Yin, Mater. Chem. Front. 4 (2020) 869e874. Copyright © 2020 Royal Society of Chemistry.

5.3.2 Applications of thermochromic perovskite materials in solar cells In the current study, the researchers were concerned about the exceptional reproducibility and the reversible thermochromic nature of dihydrated methylammonium lead iodide, distinguished by its wide bandgap that depends on temperatures ranging between room temperature and 60 C under ambient conditions because of phase transition initiated via moisture absorption and desorption. Different studies, such as X-ray diffraction (XRD) and UV-vis absorption, are performed to understand the hydratione dehydration process of the dihydrated perovskite as well as the mechanistic performance in the course of the phase transition of the reversible thermochromic materials. The thermochromic property of the obtained materials is additionally discovered in absorber materials in photovoltaic devices such as mesostructured solar cells. Additionally, the stimulated optoelectronic device features can also be used in smart windows and other buildingintegrated applications [111].

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The transition occurring in the crystallographic phases that may be related to the thermochromic behavior via XRD experiments was explored. All the prepared materials were found in the form of powder with a ratio between 1:1 and 1:4 of PbI2 to methyl ammonium iodide (MAI), which were annealed at 90 C, followed by continuous cooling to room temperature. The XRD pattern of 1:1 perovskite at room temperature confirms the tetragonal structure of MAPbI3 material, as shown in Fig. 5.6(a). Furthermore, the ratio of 1:2 as well as 1:4 of the prepared PbI2/MAI material indicates a large deviation from the original MAPbI3 perovskite structure by increasing the MAI concentration and confirming the existence of a number of new peaks (Fig. 5.6(a)). In addition, by experimentally comparing the obtained diffraction peak results for the prepared materials with the theoretical ones of the (CH3NH3)4PbI6$2H2O perovskite, a closer view could be obtained as confirmed by Leguy et al. (Fig. 5.6(b)) [112]. The crystal structure of the dehydrated perovskite is seen as a twisted NaCl structure, comprising a [PbI6] 4 octahedra and encompassed by (CH3NH3/H2O/CH3NH3)24 [113]. Besides, there is a basic uniqueness between the mono- and dihydrated MAPbI3 perovskite [114,115]. The UV-vis absorption spectra of MAPbI3 as well as (MA)4PbI6$2H2O films at room temperature and 60 C are illustrated in Fig. 5.7(a). Furthermore, from the obtained data, the band gap of the prepared MAPbI3 and (a)

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(MA)4PbI6$2H2O materials was found to be almost 1.58 and 3.82 eV, respectively [116e119]. Nevertheless, dihydrated perovskite displays a substantial alteration in the absorption spectra, with a change in color, at 60 C compared to 25 C, as shown in Fig. 5.7(aec). Moreover, one of the possible applications of (MA)4PbI6$2H2O materials are photovoltaic devices based on the reversible thermochromic property because of the phase transformation of these absorber perovskite materials. Fig. 5.7(d) demonstrates the photovoltaic performance of the device at 25 and 60 C depending of the thermochromic behavior of the absorber perovskite materials, which show a change in the performance of the assembled cells by altering the temperature parameter. The loss part issue is credited to the decline in short-circuit current density of the device, which may be due to repeated phase transition in the material. Further streamlining of the device is required to improve efficiencies. The thermochromic nature of the material is further exploited in photovoltaic devices, with uncertain power conversion efficiencies, paving the way for applications in building integrated devices.

5.3.3 Applications of hybridized VO2/graphene thermochromic materials in flexible windows The purpose of the current work is to study the construction of exceedingly crystalline and stoichiometric vanadium dioxide (VO2)-based thermochromic film by utilizing graphene on mechanically flexible substrates for the application of light transmittance and flexible smart window films able to respond to different environmental temperatures. The graphenesupported VO2 was transferred to a plastic substrate, which enabled the development of flexible VO2 thermochromic film that could be incorporated inside a mock-up house to reduce in-house temperature under infrared irradiation [120]. Fig. 5.8 illustrates a schematic diagram that shows the hybridized VO2/graphene thermochromic materials in flexible windows, supported by a scanning electron microscope photo of graphene-supported VO2 nanocomposite materials. The obtained results delivered significant advancement for the assembly of flexible thermochromic films intended for energy-saving window applications [120]. The process of large-area flexible VO2-based thermochromic film assembly is outlined in Fig. 5.9. The graphene added to VO2 is attached to Cu foil (VO2/graphene/Cu). The structure is then wrapped in a thermal release tape and floated on 0.1 M of ammonium persulfate solution to etch the Cu foils. In addition, another transfer procedure utilizing polymer

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backing film (e.g., poly(methylmethacrylate)) was also possible [121e123]. The fabrication method indicates that VO2 is deposited in a straight line on graphene-free Cu foil (VO2/Cu), and then successfully transferred to the polymer (polyethylene terephthalate [PET]) substrates. The process can confirm that graphene is a significant element to transfer VO2 to the polymer substrate. For additional identification of the morphology and structure of the prepared VO2 crystals on both graphene and graphene-free Cu substrate, a high-resolution transmission electron microscopy (HR-TEM) technique was utilized to check these samples (Fig. 5.10). Furthermore, the dispersal of polycrystalline VOx nanostructure islands, with sizes w100 nm, on the Cu substrate, which prevent the transfer to PET substrates, was detected by a cross-sectional TEM image as shown in Fig. 5.10(a). In addition, the crystal structure of the graphene-free VOx island was examined by using the fast Fourier transform (FFT) patterns, which were tested at two different spots (top and bottom region), and indicated the existence of a mixture of vanadium oxides (VOx) with a different oxidation state (VO2 and V3O5), alongside the thickness direction. Conversely, imaging of the crosssections of graphene-supported VO2 on Cu (VO2/graphene/Cu) illustrates adjoining interlinked VO2 nanocrystals on a graphene support (Fig. 5.10(b)) [124]. Besides, this information provided proof that graphene serves to help the arrangement of thickly interlinked VO2 nanocrystals and in this manner can allow the clear transfer of VO2 to PET film. In addition, the prepared graphene-supported VO2 (VO2/graphene) films have thermochromic features, which can be recognized and checked by the optical transmittance measurements of the attained temperaturedependent films as illustrated in Fig. 5.11. The fabricated films indicate typical thermochromic behavior by both optical transparency and transmission by the thermal hysteresis loop of VO2/graphene/PET temperaturedependent film as shown in Fig. 5.11(b) and (c). Despite the fact that hysteresis is significant with the multidomain VO2 construction, nanosized VO2 crystals, and other parameters [125e127], a reasonably large hysteresis recognized for VO2/graphene/PET film is moderately accredited to the existence of other VOx phases, for instance, V2O5 as well as V3O5 [128e131]. Additionally, the bending stability of the VO2/graphene/PET film was detected via checking the alteration in resistance (DR/R0) as a function of bending and strain (inset of Fig. 5.11(d)), and the obtained data revealed the joined resistance of VO2 and graphene layers [121,132].

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Figure 5.10 Structural characterization of graphene-supported VO2. (a) High-resolution transmission electron microscopy (HR-TEM) of the cross-section of as-deposited and postannealed VOx on graphene-free copper substrate (VOx/Cu). The postannealing of graphene-free VOx led to the formation of sparsely distributed VOx crystals. (b) HR-TEM of the cross-section of as-deposited and postannealed graphene-supported VO2 on copper (VO2/graphene/ Cu). The postannealing of graphene-supported VOx afforded the formation of densely interlinked VO2 crystals. The fast Fourier transform patterns after postannealing were indexed to an identical VO2 crystalline phase. Photographs show VO2/graphene laminated with thermal release tape (left) and VO2/ graphene transferred to polyethylene terephthalate (PET) film (right). Reproduced with permission from H. Kim, Y. Kim, K.S. Kim, H.Y. Jeong, A-R. Jang, S.H. Han, D.H. Yoon, K.S. Suh, H.S. Shin, T.Y. Kim, W.S. Yang, ACS Nano 7 (2013) 5769e5776. Copyright © 2013 American Chemical Society.

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The flexible VO2/graphene/PET film was used on a programmed bending machine to confirm the flexibility of the film via a continuous bending cycling test as found in Fig. 5.11(d), where the DR/R0 values of the film were preserved while it bent more than 1000 times. The results propose that the graphene-supported VO2 film can be utilized as a promising material for innovative types of energy-efficient flexible film, which directly restrict the optical transmission at certain values of temperature as shown in the prototype of the mocked-up house with VO2/graphene/PET thermochromic windows (Fig. 5.11(a)). The performance of the fabricated mocked-up house was detected via checking the alteration of the inner-house temperature after irradiation under an artificial solar lamp (Fig. 5.11(e)) [133]. By examining the data, we can confirm that the temperature is below the transition temperature of VO2 at the beginning of solar irradiation, and therefore solar irradiation can be transmitted by the two windows, giving just about the same inner-house temperature.

5.4 Applications of thermoelectric materials and devices Thermoelectric materials are considered minivoltage generators to enhance electrochemical reactions and reduce the external bias energy. Different applications of thermoelectricity have been discovered because of amplified universal energy demand linked with global warming, while waste heat organization is expected. In recent times, solar energy is seen as a renewable energy supply to power up electrochemical methods like solar cell pathways [134e149]; a temperature gradient could be naturally recognized and be used in the proposed methods. Consequently, the authors understand that the amalgamation of machine learning optimization, attached to high-throughput calculations, production, and characterization, is an essential apparatus for potential breakthroughs in thermoelectric materials. The distribution of high-throughput material combinatorial approaches is predictable to enhance process optimization rate and speed up the detection of novel materials. Accordingly, the authors mentioned trends toward a closed-loop high-throughput experimentation procedure to unravel unique physicochemical phenomena in potential thermoelectric material applications [144e149].

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5.4.1 Applications of thermoelectric silicon membrane films The authors of the current study intend to enhance the thermoelectric figure of merit (zT) in silicon membranes through a simple, low-cost, large-area, one-step experimental deposition method using an ultrathin aluminum layer. From the experiment, the deposition of aluminum can generate an ultrathin amorphous surface layer that reduces the thermal conductivity and enhances the zT value of the structure [150]. The characterization of the obtained layers via TEM indicated that the deposition of aluminum on a silicon substrate covered the surface with an ultrathin amorphous film, which showed 40%e45% higher output power as well as zT values for membrane-based thermoelectric generators. Additionally, to determine useful applications of the prepared materials, the pathway for improving the performance of a silicon membrane-based thermoelectric device and obtaining higher power generation was applied [150]. Fig. 5.12(a) and (b) illustrates the TEM images of a silicon surface before and after the sputtering deposition of aluminum via the electron beamassisted physical vapor deposition technique. Metal deposition above the flat silicon surface, with low roughness, first forms a thin amorphous layer 2e5 nm thick, followed by the polycrystalline aluminum metal [151]. Therefore through the early step of the deposition method, a thin amorphous layer can be attained as shown in Fig. 5.12(b). In addition, to detect the mechanism affected by the thermoelectric properties of silicon membranes via the deposited amorphous layer, pand n-doped samples were prepared for time-domain thermoreflectance (Fig. 5.12(c)) and four-probe electrical (Fig. 5.12(d)) measurements. The enhancement method was applied to a membrane-based thermoelectric generator (TEG) to obtain improvement in zT, which increases the net efficiency of a thermoelectric device. In-plane silicon TEG devices were assembled by connecting the hot side thermally to the substrate below, and the cold side is suspended, which is discussed by Xie et al. [152] and illustrated in Fig. 5.13(a)(c). The device contains n- and p-type silicon thermopiles, which are electrically linked in a serpentine scheme by 200 nm-thick aluminum contacts. Furthermore, Fig. 5.13(d) demonstrates the functionality of the devices as well as the open-circuit voltage, which enhanced with the hot plate temperature. Fig. 5.13(e) indicates the obtained power density at various values of temperature difference between the hot plate and room temperature (DT).

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Consequently, conversion of the ambient temperature gradient into the temperature difference across the thermoelectric membrane was enhanced via the design of the device and by creating an upper cavity structure [152]. Additionally, Fig. 5.13(d) and (e) contains data for the negative DT, where the cold and hot sides are reversed. The results obtained indicate that the quantity of power generation is low, but the device provides a quantifiable amount of thermoelectrically induced current, linearly proportional to DT. In future works, this technique is planned to be applied to materials with high zT to establish a thermoelectric device with higher power density [153e155].

5.4.2 Applications of thermoelectric silver telluride nanowire films for flexible electronic devices Ag2Te nanowires (Ag2Te NWs), as flexible thermoelectric materials, can harvest waste heat energy depending on their thermoelectric properties through the high Seebeck coefficient [156]. Although the fabrication of flexible thermoelectric films and devices paves the way to further progress in flexible electronics, the poor contact between the Ag2Te NWs results in low electrical conductivity. In the current work, the authors shared an interesting study on the effect of temperature on the electrical conductivity of the thermoelectric materials and how it was possible to fabricate Ag2Te NWs at room temperature with welding features through a simple combination of vacuum filtration and drop-casting methods. The attained materials indicate outstanding Seebeck coefficient and high electrical conductivity. Additionally, the electrical resistance of the obtained films was kept low compared with the initial one, even after 1000 bending cycles. This shows reasonable flexibility of the film with exceptional Seebeck coefficient and electrical conductivity values for a finger-touch test, as well as self-powered flexible electronic device applications via the Ag2Te NW film as a thermoelectric power generator [156]. The assembly procedure of Ag2Te NW film, where the Te NWs were produced via the hydrothermal technique and the well-dispersed Te NWs were then filtered on the filter paper, produced a substrate with good support for the obtained flexible films through the vacuum filtration pathway as illustrated in Fig. 5.14. Furthermore, the utilized vacuum filtration technique is considered a superficial and consistent method able to regulate the thickness and surface roughness of the Te NW films [157]. Te NW film added to an AgNO3 solution produced Ag2Te NW film and was able to accomplished room-temperature welding of Ag NWs.

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Figure 5.14 Schematic illustration of the fabrication of Ag2Te NW film. NW, Nanowire; PVP, polyvinylpyrrolidone. Reproduced with permission from X. Zeng, L. Ren, J. Xie, D. Mao, M. Wang, X. Zeng, G. Du, R. Sun, J.-B. Xu, C.-P. Wong, ACS Appl. Mater. Interfaces 11 (2019) 3789237900. Copyright © 2019 American Chemical Society.

Consequently, to detect and characterize the morphology of the asprepared NWs, SEM and TEM were used. The obtained Te NWs retained an average length of several microns and average diameter of 13  2 nm, with large aspect ratios of more than 300 (Fig. 5.15(a)). In addition, several welding points were detected for Ag2Te NWs (Fig. 5.15(b)). Moreover, TEM images (Fig. 5.15(c)) of the same materials affirm the efficacious synthesis of room-temperature-welded Ag2Te NW film [158]. Exhaustive microstructure examinations were approved for the Te and Ag2Te NWs (Fig. 5.15(di)). The high-resolution TEM image of Te NWs (Fig. 5.15(e)) displays strong lattice fringes, which correspond to the (100) and (111) crystal planes of the hexagonal Te (PDF# 36-1452), respectively. Also, a high-resolution TEM image of Ag2Te NWs (Fig. 5.15(h)) demonstrates a strong lattice fringe, which is in the plane of (101) for the monoclinic phase of Ag2Te (PDF #34-0142). Fig. 5.16(a) confirms the temperature-dependent electrical conductivity of the prepared welded Ag2Te NW film samples. The resulting values of electrical conductivity of the attained materials differ by changing the temperature value depending on several parameters such as the higher value of electrical conductivity, accredited to the interatomic bonding between the NWs and the thermal excitation of carriers [159]. Fig. 5.16(b) illustrates the temperature dependence of the Seebeck coefficient of two Ag2Te NW films at different temperatures, where the absolute value of the Seebeck coefficient of the two prepared films was enhanced by improving the

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temperature values (Fig. 5.16(b)), which are produced by the alteration between the average energy and the Fermi level away from the band edge. Simultaneously, the positive and negative reimbursement of the electrons and holes decreases the total Seebeck coefficient of the material, and the value of the power factor is affected. In addition, Fig. 5.16(c) illustrates the obtained optimal values for the power factors of the welded Ag2Te NW film as well as the common Ag2Te NW film [160]. Subsequently, the bending tests, by recording the changes in electrical resistance, were conducted on the welded Ag2Te NW film to check the flexibility of wearable electronic devices. The electrical resistance of the welded Ag2Te NW film and the common Ag2Te NW film is illustrated

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in Fig. 5.16(d), which shows that the room-temperature-welded Ag2Te NW film has better flexibility compared to the other film. The desirable flexibility of the Ag2Te NW film can be accredited to the large aspect ratio of the Ag2Te NWs and the joining between the NWs that can be partly eliminated in the NW network [161]. A test schematic diagram to detect the potential application of Ag2Te NW films for self-powered flexible electronic devices is shown in Fig. 5.17(a). We can notice that the voltage response can be obtained when the finger touches the Ag2Te NW film at room temperature (25 C)

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as illustrated in Fig. 5.17(b). The detected voltage should be about 0.52 mV at stable temperature, which is considered a stable output (Fig. 5.17(c)). Interestingly, when the finger touches the Ag2Te NW film, it needs only 0.4 s to attain a voltage output of 0.32 mV (Fig. 5.17(d)). For additional information on the thermoelectric performance of the film, a flexible thermoelectric module was prepared as presented in Fig. 5.17(e). The output voltage curves of the thermoelectric device can be attained under different temperature gradients by linking an external load resistance to the device (Fig. 5.17(f)). The prepared room-temperature-welded Ag2Te NW film shows exceptional thermoelectric performance that shows that these films can be used in self-powered flexible electronic devices [162e164].

5.5 Summary and perspective Progress in the optimization of physical properties and measurement techniques of thermochromic as well as thermoelectric materials paves the way for development of the thermoresponsive community. Energy harvesting has become a significant feature in the current state of progress because of the rapid breakdown of nonrenewable targets for different pathways and the increase in social mobility. In conclusion, the chapter looked at the recent different achievements used to understand improvements to different energy-harvesting thermochromic and thermoelectric devices based on several thermochromic and thermoelectric materials. Innovative models are conceived for novel approaches for the durable interrelation of thermochromic as well as thermoelectric transport parameters for thermochromic and thermoelectric devices. Additionally, this chapter offered several trials implemented in the literature to design and fabricate thermochromic and thermoelectric harvesters by highlighting their potential applications. To accumulate the efficiency and processes of equipment, the thermochromicity and thermoelectricity behaviours through thermochromic and thermoelectric materials can realized via their harvest of residual energy. The main challenges that face the fabrication process associated with both thermochromic and thermoelectric harvester implementation were discussed, as well as the considerable advantages of the targeted devices. The admission of both thermochromic and thermoelectric material harvesters in a wide range of applications, where wasted thermal energy as well as the impact of thermochromic and thermoelectric harvesters is produced, has also been highlighted.

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Technically, materials collection progress can be enhanced via highthroughput computational systems to achieve preliminary screenings of the search area and recognize the best applicants to be considered experimentally, namely, the so-called materials genome approach. Materials with intrinsically very low thermal conductivity are also of great interest, especially when they simultaneously feature unique band structures (for example, the conducting bands are composed of multicarrier pockets) that allow high power factors. Finally, in view of the complexity of both thermochromic and thermoelectric materials, their future growth would benefit from collaborations between chemists, physicists, and materials scientists. In addition, the requirement of substitute energy technologies and materials to support the reduction of the use of fossil fuels to fulfill the need of environmentally friendlier technologies would allow thermochromic and thermoelectric energy conversion technology to play an important role in future research and applications.

Acknowledgments AES is currently on leave from CMRDI. In addition, BCMaterials is thanked by the authors for their help in following up this study. Furthermore, AES thanks the National Research grants from MINECO “Juan de la Cierva” [FJCI-2018-037717].

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CHAPTER SIX

Lightweight, multifunctional materials based on magnetic shape memory alloys Daniel Salazar-Jaramillo1, Jose M. Barandiaran1, 2 Manfred Kohl3, Daoyong Cong4, Hideki Hosoda5, Jose Luis Sanchez Llamazares6, Volodymyr A. Chernenko1, 2, 5, 7 1

BCMaterials, Basque Center for Materials, Applications and Nanostructures, Leioa, Spain Department of Electricity & Electronics, University of the Basque Country, Bilbao, Spain 3 Karlsruhe Institute of Technology, IMT, Karlsruhe, Germany 4 Beijing Advanced Innovation Center for Materials Genome Engineering, State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing, China 5 Institute of Innovative Research (IIR), Tokyo Institute of Technology, Yokohama, Japan 6 Instituto Potosino de Investigaci
on Científica y Tecnol
ogica A.C., San Luis Potosí, Mexico 7 Ikerbasque, Basque Foundation for Science, Bilbao, Spain 2

6.1 Introduction Contemporary global challenges in efficient energy consumption and material resources use need multifunctional materials enabling new technologies. Multifunctional magnetic shape memory alloys (MSMAs) can meet these challenges. They usually comprise those off-stoichiometric Heusler-type X2YZ compounds (where X and Y are transition metals and Z belongs to the IIIeV group in the periodic table of elements), which exhibit thermoelastic martensitic transformation (MT) accompanied by a complex magnetic reordering. The discovery and start of research of the prototype of MSMAs, namely NiMn-Ga alloys family, date back to the early 1990s [1e3]. Due to strong interactions between the structural and magnetic degrees of freedom in the vicinity of MT, MSMAs may exhibit a large change in some properties that define their functional giant properties, such as magnetic field-induced strain (MFIS), magnetoresistance, direct and inverse magnetocaloric effect, alongside the pronounced conventional shape memory effect and superelasticity. These properties can be readily tuned by composition variation, doping, or heat treatments. MSMAs are cost-effective, rare-earth-element free, and nontoxic. They can be conventionally divided into two major groups: (i) ferromagnetic shape memory

Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00005-6

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alloys (FSMAs), such as aforementioned prototype Ni-Mn-Ga, especially suitable for actuation, sensing, damping, and vibration energy harvesting owing to the unique reorientation effects of a martensitic structure; and (ii) metamagnetic shape memory alloys (meta-MSMAs), such as Ni-Mn-X (X ¼ In, Sn, Sb), which are especially important for ferroic solid-state cooling as a result of stress- or magnetic fieldeinduced MT. Both types of materials are capable of efficiently transducing the thermal, mechanical, and magnetic energies into each other under conjugating fields giving rise to the cross-linked new physical effects on both macro- and micro(nano)-scale. FSMAs offer various options for actuation. One example is the effect of magnetic fieldeinduced twin boundary motion in the martensitic phase being the underlying mechanism of the magnetic shape memory (MSM) effect [3,4]. Ni-Mn-Ga FSMAs exhibit strong magneto-elastic interactions and low twinning stress. Thus, an external magnetic field exerts an equivalent magnetostress which is enough to move twin boundaries causing an increase in the fraction of magnetically favorable martensite variants at the expense of less favorable variants. This magnetically induced reorientation process is accompanied by a giant macroscopic strain that depends on lattice parameters. In the case of Ni-Mn-Ga exhibiting a 10M modulated tetragonal martensite, the reversible MFIS reaches 6% [5]. Depending on twin morphology, Ni-Mn-Ga may exhibit an extremely low twinning stress, down to less than 0.1 MPa, which has to be overcome by the magnetostress to enable reorientation of the martensite variants [6]. In this case, moderate external magnetic fields, even below 0.1 T, are sufficient to induce the MSM effect. The MSM effect has been shown to enable fast actuation with the twin boundary motion speed in the order of 0.5 m/s [7]. By surface treatments, an increased number of pinning sites are established that prevent the coarsening of twins and thus enable the formation of a stable twin microstructure with mobile thin twins, which is the basis for high performance MSM actuation [8]. The MT in MSMAs, involving large abrupt change in both structural and magnetic properties, offers actuation by the thermal shape memory effect (SME), the thermally induced change of magnetic force (TMF), and the giant (magneto)elastocaloric response in such materials. The multifunctional properties and high power density, the latter one resulting in an excellent scaling behavior of mechanical properties, render MSMAs to be particularly attractive for small-scale applications [9]. Thus, the actuation capacity of MSMAs arises from their intrinsic (thermo)magnetomechanical properties, allowing for a drastic reduction in size, weight, and complexity of the devices, that is in line with a general

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demand for the magnetic functional materials of 21st century to be stronger, lighter, and more energy efficient [10]. This chapter presents a brief overview of the MSMA materials in the lightweight low-dimensional form, such as the powder/polymer composites, foams, thin films, micropillars, thin wires and melt-spun ribbons, exhibiting an excellent functionality on the microscopic scale.

6.2 FSMA/Polymer composites for magnetic actuation and mechanical damping 6.2.1 Magnetic fieldeinduced rubberlike behavior of Ni-Mn-Ga/silicone composite The motivation for the development of the magnetostrain active MSMAs/polymer composites stems from (i) some disadvantages of the bulk single crystalline MSM actuators, such as their high fabrication costs, enhanced brittleness, and frequency limitations, and (ii) incapacity of the polycrystalline bulk MSMAs to exhibit a large enough magnetostrain due to constraints from grain boundaries (see Refs. [11e16] and references therein). The obvious key precondition for a design of the magnetostrain active Ni-MnGa/polymer composites, suitable for the magnetic fielde induced actuation via field-induced twin boundaries motions in the embedded particles, is that the individual particles should be in the martensitic state and exhibit a similar twin structure mobility as in a bulk single crystal, which means that their twinning stress should be lower than the magnetostress (about 1e3 MPa at a saturating magnetic field below 1 T). Furthermore, a stiffness-matched-matrix, i.e., the polymer with the sufficiently low stiffness allowing a field-induced shape change of the particles and optimal amount of particles assembled in the chains are needed to expect a substantial MFIS of the whole composite and its reversibility [16e18]. In this case, the embedded particles should be crystallographically oriented providing that the easy-magnetization short axes of tetragonal unit cells are aligned along the chains. The optimal volume fraction of particles is selected to avoid a dilution effect on the MFIS characteristics of composite being compromised with the upper bound limit of volume fraction, where the deformation of particles is blocked by the neighboring particles [16,19]. The nonfulfillment of, at least, one of the aforementioned conditions results in a Ni-Mn-Ga/ polymer composites exhibiting less than 0.1% of MFIS [20,21]. All the above mentioned conditions have been realized for the first time in Ref. [16]. Fig. 6.1 shows a large value of magnetostrain, 4% in elongation

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4

Elongation Magnetostrain [%]

2

H

0

–2

Contraction –4 0

0.5

1

1.5

Magnetic field [T]

Figure 6.1 Magnetic fieldeinduced deformation of Ni-Mn-Ga particles/silicone composite. The solid curves show a sample contraction along the magnetic field, H, accompanied by its elongation in the orthogonal direction. Dashed curves depict the strain recovery as the magnetic field removes. Adapted from P. Sratong-on, V.A. Chernenko, J. Feuchtwanger, H. Hosoda, Magnetic field-induced rubber-like behavior in Ni-Mn-Ga particles/polymer composite, Sci. Rep. 9 (2019) 3443, https://doi.org/10.1038/ s41598019-40189-2.

and about 2% in contraction, of the Ni-Mn-Ga particles/silicone composite sample under applied field of less than 0.7 T. Composite starts to deform when the magnetic field reaches the value, about 0.3 T, necessary to move twin boundaries in the particles. The complete shape recovery of the sample during field removal is observed with some hysteresis. Four percent of MFIS in composite is comparable with a maximum achievable magnetostrain in the Ni-Mn-Ga bulk single crystal (6%) at the same orientation between applied field and appropriate crystallographic direction [5]. Also the value of the switching field for activation of the twin boundaries motion, and the field value when they stop movement, equal to 0.3 and 0.7 T, respectively, are typical for the bulk single crystal [5]. The principal difference in the MFIS behavior between the bulk Ni-Mn-Ga and composite consists in the fundamental fact that, in the bulk samples, the magnetostrain recovery is only possible when the external orthogonal force is applied. Indeed, by whatever force a martensite is deformed through twinning mechanism, there is no thermodynamic driving force to reset its initial shape. So, a large inherently recoverable MFIS, shown in Fig. 6.1, is recognized as a new effect called as the magnetic

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Figure 6.2 X-ray mCT 3D-image of 30 vol.% NiMnGa/silicone composite. The bold arrow indicates a direction of the particle chains. Adapted from P. Sratong-on, V.A. Chernenko, J. Feuchtwanger, H. Hosoda, Magnetic field-induced rubber-like behavior in Ni-Mn-Ga particles/polymer composite, Sci. Rep. 9 (2019) 3443, https://doi.org/10.1038/s4159801940189-2.

field-induced rubberlike behavior of MSMA/polymer composite. In the case of composite, the backward twin boundaries movements during field removal were shown experimentally and theoretically to be produced by the influence of local internal stresses [16,19,22]. During field application, the particles are deformed creating biasing stress fields in the surrounding polymer which serve as a driving force for the subsequent strain recovery. It was shown that the significant biasing stress fields appear if the particles are elastically interacting. Particularly, this happens when they are located at a distance less than 500 mm from each other, creating pairs or groups of the nearest neighbors [22]. Fig. 6.2 illustrates the X-ray mCT 3D-image of 30 vol.% NiMnGa/silicone composite which was used in the magnetostrain measurements (Fig. 6.1). The particles are aligned along the chains during curing process of composite under magnetic field. The curing is performed in the martensitic state facilitating the selection of twin variants oriented with their short easymagnetization c-axis along the particles chains. As a result, an artificial magnetic anisotropy is introduced to the whole composite, as can be confirmed by the magnetization curves measurements. The magnetostrain results, depicted in Fig. 6.1, are obtained under application of the magnetic field perpendicularly to the chains. The main key factor causing such a giant

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MFIS of composite lies in the magnetic fieldeinduced high mobility of twin boundaries in the Ni-Mn-Ga particles resulting in the reorientation of twin variants with their short c-axis aligned in the field direction. The twin rearrangements become possible due the special-quality particles used [23] which are shown in Fig. 6.3. They represent the individual single crystalline grains faceted by the atomic planes with a minimum of the crystal defects. One type of twinning planes spreading across a particle, as can be appreciated in the inset to Fig. 6.3, evidences a rather good single crystallinity of particles. Such particles are prepared following the innovative procedure described in Ref. [16]. This procedure includes a know-how consisting in doping of an alloy, such as Ni49.4Mn28.6Ga22.0 (at.%) exhibiting MT near 305K, with a tiny amount of Bi (0.03 at.%). Bi element fully segregates on the grain boundaries, thereby radically increasing an intergranular brittleness. After careful mechanical cleaving of the polycrystal and subsequent heat treatment, the obtained powder consists of the already mentioned undeformed single crystalline virgin grains, which exhibit the one-way magnetic fielde induced twin boundaries displacement in a free state, as directly observed in Ref. [16]. The X-ray mCT 3D-imaging revealed that this twin boundary displacement, generally responsible for the large shape change of any FSMA sample, appears reversible in the grain particles distributed inside of the silicon polymer composite, whereby providing a crucial support for

200 Pm

Figure 6.3 SEM image of single grains separated by a gentle disintegration of the Ni49.4Mn28.6Ga22.0-Bi ingot. Inset depicts a 3.5-magnified image of one particle providing better view of twin structure exhibiting a single orientation of the twin boundaries. See also Ref. [23].

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the explanation of a large recoverable MFIS of the whole composite and magnetic fieldeinduced rubberlike effect. A large MFIS of the selected individual grains inside of composite was directly determined by the analysis of the three-dimensional mCT images obtained in-situ before application of the magnetic field, under magnetic field and after its removal [16]. In addition, one should mention an important role of the silicone matrix in getting large reversible MFIS of composite. In this case, a silicone rubber ELASTOSIL M4400 with a stiffness of 0.5 MPa ensured a stiffness-matched medium enabling both the particles deformation and accumulation/release of the backwards internal stress in the matrix during application/removal of the magnetic field. A stress release occurs to restore an original shape of the particles. The usage of an FSMA/polymer composite material exhibiting a magnetic fieldeinduced rubberlike effect can significantly simplify the actuation or sensing devices, since no external biasing force is required to reset it after removal of the magnetic field.

6.2.2 MSMA composites for magnetoelectric applications Magnetoelectric (ME) composite materials, consisting of magnetostrictive and piezoelecric components, produce electrical signal under a magnetic field application (direct ME effect), and/or exhibit a magnetization change when subjected to an electrical field (converse ME effect) [24]. The elastic interactions across the interfaces between the mentioned components control an efficiency of ME coupling. The ME effect offers novel promising applications in magnetoelectric energy harvesters, actuators, sensors, memory storage devices, etc. [25e27]. A great interest in the literature is given to the development and studies of laminated ME structures, where the giant magnetostrictive materials are used, such as TERFENOL-D (see Ref. [24] and references therein). As far as FSMAs have almost one order of magnitude larger magnetostrain response than TERFENOL-D, they have been already tried as a magnetostrictive component in ME laminated composites with PZT [28], with PMN-PT [29,30], and the PVDF piezopolymer [31]. The converse ME properties of the sandwichlike layered composites consisting of a PMNPT single crystal plate bonded between the two Ni-Mn-Ga single crystal plates have been studied in Ref. [29]. A giant resonance value of the converse ME coefficient, equal to 18.6 G/V, was measured at the MT temperature under a very low bias field. The mechanism of such an effect is associated with the strain-induced twin variant rearrangements, resulting in the considerable change in magnetization of Ni-Mn-Ga [32].

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A sandwichlike composite consisting of one layer of Ni-Mn-Ga single crystal plate bonded between two layers of the piezoelectric PVDF polymer film was fabricated, and its MFIS and ME characteristics were studied as a function of magnetic field and mechanical load in Ref. [31]. The composite has demonstrated MFIS of 0.35%, whereas the bulk Ni-Mn-Ga, which was used to prepare thin plate, has shown MFIS of 5.6% in the calibration tests. Such a difference in the magnetostrain values is explained by the constraint conditions of FSMA in the composite. The measured ME coefficient in the external loadfree condition was equal to 0.58 V/cm$Oe under a magnetic bias field of 8.35 kOe, which was attributed to the twin-rearrangementorigin of MFIS in FSMA layer. This result motivates the future efforts in development of the Ni-Mn-Ga/piezopolymer composites exhibiting a large ME effect. To this end, PVDF thin films containing a system of magnetostrain active particles, such as described in Section 6.2.1, would be very promising.

6.2.3 FSMA/polymer composites for mechanical damping Sections 6.2.1 and 6.2.2 demonstrate that the elaboration and research of FSMA/polymer composites for the magnetic field actuation and magnetoelectric energy conversion are still in their infancy. More matured is a research field of the Ni-Mn-Ga/polymer composites for mechanical damping, to which the strain-induced twinning in the Ni-Mn-Ga particles is a major contribution. Indeed, the twin boundary motion in the particles, as the main mechanism for a large absorption of the mechanical energy, can be easily activated by mechanical stress [20,23,33,34]. Therefore, to design FSMA/polymer composites for mechanical damping, the magnetostrain active particles, described in Section 6.2.1, are desirable but not strictly necessary, that is why in this case FSMA powder is usually obtained by the traditional powder metallurgy methods, such as ordinary crashing, ball milling, gas atomization, or spark-erosion. In addition, polymers with an enhanced stiffness such as the polyurethane or epoxy resin can be used. FSMA/polymer composites for mechanical damping are especially interesting because of the potential to be controlled by the magnetic field if the appropriate stress-assisted conditions are realized [34]. Some advances in this research field, extracted from 14 publications, have been summarized previously (see Ref. [13] and references therein). Recently, FSMA/polymer composites have been considered for vibration [14,23,34] and impact damping [35]. The stress-induced twin boundary motion, as the main mechanism of damping, was confirmed directly or indirectly. This mechanism is

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largely dependent on the value of stress amplitude applied to a composite. Both low and high values of stress amplitude were explored by a dynamic mechanical analysis (DMA) performed as a function of temperature in a single cantilever [34] or uniaxial vibration loading [23] modes. Particularly, it was demonstrated that the gas atomized Ni-Mn-Ga powder/epoxy composite has a stiffening effect under applied magnetic field which is comparable to that one of the single crystalline Ni-Mn-Ga [34]. Furthermore, it was found that in this composite the magnitude of the magnetic fieldeinduced stiffening is directly related to the filling ratio. The MSMA/polymer composites are promising for damping applications, as they are easy to fabricate and their mechanical properties are easy to vary. Inasmuch as the vibration energy absorption of these composites can be controlled by an external magnetic field, the novel damping applications could be anticipated.

6.3 Porous FSMA materials One approach to eliminate an influence of the intergranular constraints on the twin boundaries motion under a magnetic field or an external mechanical stress is to separate an FSMA polycrystal into individual grains and then embed them into polymer to create magnetostrain active textured composite, as has been described in Section 6.2.1. To cleave polycrystal into a single crystalline grain powder, a special technology, briefly described in Section 6.2.1, was elaborated. The other approach to drastically reduce an influence of the grain boundaries constraints is a partial removal of some amount of grains to create a porous material with an irregular strut-linked-by-nodes architecture. This idea was successfully realized by the elaboration of processing of the Ni-Mn-Ga foams [12,36,37]. The foams were shown to exhibit a giant magnetic-field-induced strain, up to 9%, in a cyclic regime provided by the rotating magnetic field of about 1T [12]. This outstanding value of MFIS is larger than the one expected from crystallography, which means that some contribution from the magnetic fieldeinduced torque takes place. In fact, with one types of pores, having size comparable to the size of the grains, only 0.12% of MFIS was observed, but after additional introducing the system of much smaller pores and heating/cooling through MT, MFIS of the 62% porous sample reached 8.7% [12]. The microscopic analysis has shown oligocrystalline structure, where struts represent single crystalline areas which are linked by the nodes dotted by a multitude of small pores. The latter ones help to further reduce

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constraints which hinder twin boundaries motion in the struts. Several series of the cyclic magnetomechanical tests have shown that it is possible to get a stable MFIS signal in more than 200,000 cycles. MFIS value can be varied within 2.0%e8.7%, depending on the sample history [12]. Whereas the Ni-Mn-Ga foam in Ref. [12] was created by the replication method, using liquid metal infiltration into a preform porous ceramic, in Ref. [36] the foam was produced by the spark plasma sintering of Ni-MnGa and NaCl powders. The latter approach was effective to easily control the porosity of the Ni-Mn-Ga foam, between 60% and 90%, thereby obtaining MFIS value versus porosity dependencies. In this case, the highest MFIS of 1.24% was achieved on the sample with the highest porosity. The effect of directional solidification on texture and MFIS in Ni-Mn-Ga foams with 57 vol.% of 355e500 mm open pores and coarse grains was studied in Ref. [37], where a maximum cyclic magnetostrain of 0.65% was measured. Due to quite complicated technologies of synthesis of Ni-Mn-Ga foam, hard possibility to reproduce a desired foam architecture and related functionality, the modern trend now is replaced by the different methods of additive manufacturing which still are operating with the bulk MSMA samples as the model systems.

6.4 MSMA thin films and related nanostructures 6.4.1 Development of thin films Most of the modern functional materials do not need to be used in bulk as they can display their properties in very small sizes, such as small particles, thin films, thin wires, or micropillars. Thin films are defined as material layers having less or about 1 mm thickness. Compared to bulk materials, [38] the area-to-volume ratio is very large in thin films, which can alter the mechanical, thermal, electrical, or magnetic properties of the material. Therefore, when a material is reduced to the nanoscale, new applications for devices based on micro or nanostructured materials are possible. MSMA thin films and nanostructures are of special interest because of their high power-to-weight ratio and the large strains they can undergo under heating/cooling, stress, or magnetic field. These features make them promising materials for MEMS [39], where they could play as actuators or sensors. In most cases, MSMA thin films are deposited onto a substrate that serves only as a support. Free-standing films can be obtained by depositing onto

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some substrates with a subsequent peeling off. Alternatively, they have been fabricated using a sacrificial layer, which is removed by chemical etching [40] or by a sacrificial substrate like NaCl, afterward dissolved in water [41]. Freestanding cantilevers and bridges of quite large dimensions (100e400 mm) have been sculpted from a single crystal by photolithography and reactive ion etching [42]. A short introduction to the development of MSMA thin films is provided below. A more detailed description can be found in [43]. Thin films can be prepared by a number of physical and chemical methods, although not all of them are suitable for metals and alloys. By far, the most used methods for depositing MSMAs are the Physical Vapor Deposition (PVD) ones, such as sputtering, pulsed laser deposition [44,45], ion beam deposition [46], flash evaporation [47] and molecular beam epitaxy [42]. Magnetron sputtering is a widely used technique that conserves rather constant the composition of the target alloy in the deposited films [48]. However, the properties of MSMAs are very sensitive to their composition; so, in addition to the target composition, all the parameters influencing the deposition, such as gas pressure, deposition rate, substrate type, and temperature, target to substrate distance, etc., must to be carefully controlled. For instance, the Ni content was found to decrease by increasing the sputtering power [49] and the residual stress in the film is very sensitive both to the Ar pressure and deposition power [50]. The nature of the substrate has a strong influence on the film structure. For instance, a single crystal with lattice parameters as close as possible to those of the film promotes epitaxial growth, and the resulting film is also a single crystal. MgO [51], Al2O3 [52,53], and SrTiO3 [54] single crystals have been used as substrates to grow epitaxial Ni-MnGa FSMA films. As discussed below, different film orientations can be obtained on MgO (001) depending on the thickness of the film [55]. Polycrystalline and amorphous substrates, however, give rise to a nanocrystalline granular growth either fully isotropic or partially textured. Si (001) with a naturally or intentionally grown layer of SiO2 or nitride surface layer is the most representative [45,50]. Other polycrystalline substrates are alumina [56], mica [47], or molybdenum [57]. Elastically soft substrates, like polyvinyl alcohol [58] or photoresist [59] allow to grow thick films, up to 10 mm, without auto peeling off, because the substrate absorbs the stress and has an enhanced adherence. The substrate temperature is critical to get the desired compositional homogeneity, degree of crystallinity, crystal structure, and texture of

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Ni-Mn-Ga films. Careless depositions result in samples with a low crystallization degree (nanocrystalline or even amorphous) showing neither ferromagnetism nor martensitic transformation. Only high temperature deposition (500 C or higher) [43,60,61] enables the desired crystal structure and homogeneity of the films. Residual stresses in the film are extremely important because they can modify the martensitic transformation temperature, twin variants selection, and magnetic behavior. As the deposition or thermal treatments take place usually at elevated temperatures, cooling down induces a residual stress in the film due to the different thermal expansion coefficients of film and substrate. Stress in the film can be estimated through the expression [43]: E ðaf  as ÞDT (6.1) 1v Here DT ¼ (T0T), T0 is the film deposition temperature, or the temperature of an annealing treatment, at which the stress is assumed to be zero; E and v are the Young’s modulus and Poisson ratio of the film, and af and as are thermal expansion coefficients of film and substrate. In most Ni-Mn-Ga thin films, the substrate produces a positive Da ¼ afas resulting in the in-plane tensile stresses in the film. NaCl, however, has a thermal expansion coefficient larger than that of Ni-MnGa, so that Da 10 6

K/s

Grain size (Pm)

Molten alloy Coo

(c)

Thickness (Pm)

Induction coil

Nozzle

(b)

(a)

Wheel speed (ms–1)

ed

Powder Ribbons Flakes

Figure 6.16 Scheme of melt-spinning technique.

typically varied between 5 and 60 m/s. Whereas the amorphous materials can be obtained as the kms-long ribbons, in the case of materials with a fast crystallization and growth kinetics the ribbon pieces with lengths of 3e150 mm can be only expected. Although the rapid solidification induces some degree of disorder and stress depending on the cooling rate [145], a subsequent heat treatment at moderate temperatures, e.g., around 1073K, can be performed to homogenize the phases and to improve the chemical and crystalline order. From a manufacturing point of view, this technique allows to obtain large amounts of material in a short time that ensures a large-scale production of MSMA ribbons or MSMA precursor material for any of their potential applications. Another feature observed in MSMA ribbons is the preferential crystallographic orientation resulting from the crystallization conditions of the processing method [144,146e148]. As it is known, the recoverable strain exhibited by the shape memory alloy is directly linked to the crystal orientation, being enhanced when the deformation occurs along the texturing direction. On the other hand, since thin melt-spun ribbons of MSMAs are quite brittle, they can be utilized for the simple production of fine powders with predetermined physical properties avoiding a severe milling process that could induce significant surface microstrains, defects, and high disorder in the milled particles or contamination.

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This section focuses on the recent progress in meta-MSMA thin ribbons with emphasis on their functional properties, such as magnetoresistance and caloric effects.

6.6.2 Martensitic transformation The main ingredient of MSMAs is the first-order diffusionless MT. As already mentioned, MT in meta-MSMA is accompanied by a change of the magnetic state in the material. This change implies that the lowtemperature martensitic phase is weakly magnetic, where the magnetic spin configuration could be antiferromagnetic, superparamagnetic, ferrimagnetic etc., whereas the high temperature austenitic phase is normally ferromagnetic. As MT proceeds, e.g., on heating, martensitic phase progressively transforms into austenite, which can be not only ferromagnetic but sometimes also paramagnetic depending on the position of Curie temperature of austenite with regard to the MT temperature. Thus, complex phase diagrams can be realized in meta-MSMAs offering cross-related properties that may lead to important applications in the magnetic refrigerators, sensors, or actuators. It is noteworthy, that the quasi-equilibrium MT phase diagram is strongly dependent of the microstructure of the alloy. Taking into account that meta-MSMA melt-spun ribbons have been obtained without additional heat treatments [149,150] and that first-order MT, generally, proceeds through the mechanism of nucleation at dislocations, grain boundaries or other defects, the effect of solidification rate on MT and on the structural properties has been studied in as-spun ribbons of Ni50Mn37Sn13 (v was varied between 15 and 50 m/s) [151] and Ni50Mn35In15 (v was varied between 10 and 50 m/s) [152]. For Ni50Mn35In15 ribbons it was reported that MT shifts to lower temperatures at a rate of 0.4 ks/m with the increasing of the linear speed of the Cu wheel; in addition, austenite crystallizes in a mixture of the L21 and B2 structures for 10  v  30 m/s and in the pure B2-ordered phase for higher wheel speeds. In all the cases, the reduction of MT temperature was correlated with an effect of the average grain size refinement.

6.6.3 Magnetoresistance The brittleness of MSMA ribbons makes challenging the realization of electrical measurements using the four-probe method and, therefore, the assessment of magnetoresistance (MR) around MT. The difficulties with measurements can explain only a few number of works reporting on the

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MR effect associated to MT in MSMA melt-spun ribbons. Recently, special attention has been paid to Ni-Mn-Sn-based meta-MSMAs due to their potential for the elastocaloric and sensor applications [149]. MR and exchange bias effects have been studied in the as-solidified Ni41Mn50Sn9 melt-spun ribbons obtained at 20 m/s [153]. The value of MR at 10K was measured to be quite low, 0.8% for m0H ¼ 5T, whereas a large exchange bias field of w1.13 kOe was observed at this temperature. There are other works mainly focused on the MR characterization. Electrical properties of the Ni43.5Mn44.7Sn11.8 meta-MSMA melt-spun ribbons, annealed at 1073K for 1 h, have been studied [154]. According to XRD measurements, the crystal structure of these ribbons is a single austenite phase at room temperature. The microstructure has a columnar grain character with a preferred [400] crystallographic orientation of the crystallites perpendicularly to the ribbon plane. Such a microstructure caused a two times larger MR if compared with a bulk parent sample. A large MR in the ribbon plane, equal to about 25%, was obtained at 276 K due to a magnetic fieldinduced MT, when a magnetic field of 1T was applied perpendicularly to the ribbon plane [154]. The Mn loss by evaporation during the fabrication process of the meta-MSMA ribbons is a usual problem that must be overcome to get the desired stoichiometry. It was reported that the evaporation of Mn could be prevented by adding a small amount of boron to the alloy [155]. The same study also reports the magnetotransport properties of Ni48.95Mn37.74Sn13.32B3 ribbons exhibiting MT around room temperature. A maximum MR value of 32% was obtained for a magnetic field of 9 T. An anomalous Hall effect around MT was also reported, the highest voltage change in a magnetic field of 9 T at 100 mA was 1.2 mV [155], which is suitable for magnetic sensor applications. The same group studied the effect of Fe doping on MR in Ni50-xFexMn38Sn12þB3 (x ¼ 0, 1, 3, 5) ribbons [156]. MT decreases with Fe doping from 290 to 213 K, thermal hysteresis increases from 16 to 21.5 K, and the absolute value of the magnetic fieldeinduced shift of the MT temperature increases from 7 to 35 K when a magnetic field of 5T is applied. Maximum MR value of 50% at 204 K for a magnetic field of 9T was reported for the sample with x ¼ 5 [156]. Recently, MR of the Ni50-xCoxMn35Ti15 (x ¼ 13 and 13.5) meta-MSMA ribbons produced at 15 m/s and heat treated at 650e850 C under Ar atmosphere for 10 min has been reported [157]. MT shifts by 170 and 80 K to higher temperatures after annealing for x ¼ 13 and 13.5, respectively. The MR measurements performed in these annealed ribbons

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by applying a field of 5T gave values of 26.5% at 336 K and 34.9% at 295 K for x ¼ 13 and 13.5, respectively [157].

6.6.4 Magnetocaloric effects in MSMAs Discovery of a giant MCE associated to first-order transition and its relevance from the viewpoint of solid-state refrigeration applications has catalyzed investigations of the magnetocaloric properties of (Ni, Mn)based MSMA systems prepared both in bulk and by the rapid solidification in the form of melt-spun ribbons. Hereafter, we shall only focus on the most relevant results reported on melt-spun thin ribbons. The assessment of MCE properties of rapidly solidified Ni-Mn-Ga meltspun ribbons was made for Ni55Mn20.6Ga24.4 and Ni55Mn19.6Ga25.4 alloys thermally annealed at 1075K for 3 h [158]. In these alloys a ferromagnetic martensite transforms into a paramagnetic austenite around room temperature with relatively low thermal hysteresis of 6 and 9K, respectively, giving rise to a narrow thermal dependence of the magnetic entropy change DSm(T). For a magnetic field change of 2 T, these alloys show maximum absolute values of the magnetic entropy change.jDSm jmax of 9.5 and 10.4 J/kgK, respectively, around 309K. Magnetic hysteresis loss across the phase transition was found to be negligible. An essentially larger DSm around 354 K was reported for Ni52Mn26Ga22 ribbons thermally annealed at 1173K for 18 h [159]. The authors emphasized that the annealing considerably enhanced the atomic ordering degree and suggested that an intermartensitic transformation coupled with the magnetostructural transformation occurs because of the co-appearance of a mixture of monoclinic 7M- and nonmodulated (NM) martensitic phases. The strong coupling between the magnetic transition and the martensitic (intermartensitic) transformations led to jDSm jmax values of 16.4 J/kgK and 30.0 J/kgK for 2 and 5 T, respectively. The first studies about the MCE response associated to MT in rapidly solidified ribbons of (Ni-Mn-X)-based meta-MSMAs with X ¼ Sn and In, were carried out for Ni50.3Mn35.5Sn14.4 [144] and Mn50Ni40In10 [160] Heusler alloys. In both cases, the moderate absolute values of jDSm jmax were obtained owing to the reason that MCE was not high for these compositions. The occurrence of a large hysteresis loss associated to the magnetic fieldeinduced MT in the ribbons was emphasized. The conclusion has been made that in a textured Mn50Ni40In10 ribbons jDSm jmax , magnetic hysteretic loss and refrigerant capacity are not significantly affected by the texture [161]. Nevertheless, these works stimulated further studies on

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MCE in rapidly solidified melt-spun ribbons of these families of alloys. A giant jDSm jmax of about 34 J/kgK for a magnetic field of 5T was reported at 300K for Ni45Mn37In13Co5 ribbons thermally annealed at 973K for 24 h. The samples were characterized by a wide thermal hysteresis of MT (of 18K at 10 mT that increases with the increase of the applied magnetic field) and large magnetic fieldeinduced hysteresis loss [161]. A large magnetic entropy change associated with a high-temperature martensitic transformation, from a paramagnetic-type six-layered modulated martensite to a ferromagnetic austenite, occurring around 395e380 K has been recently reported for Ni40Co10Mn41Sn9 melt-spun ribbons [162]. For a magnetic field change of 5 T (2 T), the magnetic entropy changes of 33.9 (25.0) and 53.0 (22.0) J/kgK were measured through the reverse and forward martensitic transformations, respectively. These values are among the largest reported so far for this family of Ni-Mn-based Heusler alloys. As in all alloys of this family, a large magnetic hysteresis loss associated to effect of the applied magnetic field on the reverse martensitic transformation was detected. Recently, MT from the antiferromagnetic martensite to the ferromagnetic austenite, has been observed in the all-d-metal quaternary Ni-Co-Mn-Ti Heusler alloy system [163,164]. By properly choosing the composition, the MT temperature can be tuned in a wide temperature range including room temperature. Thermally annealed melt-spun ribbons in this alloy system show a remarkable enhancement of their magnetocaloric response in comparison with the observed ones in their bulk counterparts; this has been attributed to the high degree of chemical homogeneity that the samples obtained by this method show [165]. For a magnetic field change of 2T, Ni37.5Co12.5Mn35Ti15 ribbons, annealed at 1073K during 30 min, show at 290 K jDSm jmax of about 27 J/kgK, which is among the largest values reported for most of the room temperature magnetocaloric materials [165].

6.7 Conclusions Recently, considerable efforts have been undertaken to fabricate and investigate low dimensionally shaped Heusler-type MSMAs. As a result, new light-weight multifunctional materials, such as MSMA/polymer composites, foams, thin films, micropillars, thin wires, and ribbons have been developed for small-scale applications as actuators, sensors, and energy conversion devices. A brief overview of the recent results in this area of

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MSMA research and technology is given in the present chapter and the following conclusions can be outlined. The specially designed magnetostrain active polymers containing Ni-Mn-Ga particles were shown to exhibit a large magnetic fieldeinduced rubberlike effect. Further studies are needed to demonstrate that such polymer composites have an unbeatable functionality, e.g., in the haptic applications or actuators. Besides, such polymers are very promising for magnetic fieldecontrolled damping applications and for the vibrating energy harvesting. Two opposing thermal actuation principles in a single FSMA device have been demonstrated in the Ni-Mn-Ga thin foils and films resulting in the application as the bidirectional microactuators for optical scanners [166]. Recent work has shown that the Ni-Mn-Ga thin foils and micronanopillars exhibit magnetic shape memory effect opening up fast magnetic-field-driven actuation at ultra-small scales. Free-standing FSMA film and bimorph nanoactuation have a huge potential for photonics applications, e.g., optical nanoswitching [167] and directional coupling [168]. Magnetic shape memory microwires show a well-pronounced temperature/stress/magnetic fieldeinduced MT, whereby the giant tensile superelasticity, shape memory effect, and magnetocaloric effect have been achieved. Since microwires have advantages, such as a high specific surface area and enhanced mechanical properties, they can be directly used in miniature sensor, actuator, and solid-state refrigeration devices, or act as the building blocks for macro-devices. More efforts should be undertaken in the technological aspect of the MSMA thin wires, in order to optimize their transformation characteristics and to improve their reproducibility and cyclic stability. MSMA melt-spun ribbons can exhibit a very sharp MT characteristic under applied temperature or magnetic field evidencing their good homogeneity. Under external stress they become more brittle than the thin wires, reflecting their polycrystalline nature compared to the oligocrystalline structure in wires. Therefore, meta-MSMA ribbons, as having also a potential of a large-scale production, are more suitable for the magnetocaloric and/or magnetoresistance applications. While the melt-spinning technique can provide good quality and high surface/volume ratio ribbons without additional heat treatments, still big efforts should be made for controlling the fabrication procedure to obtain MSMA ribbons with the highly reproducible functionalities.

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Acknowledgments The financial support from Spanish Ministry of Science, Innovation and Universities (RTI2018-094683-B-C5), National Natural Science Foundation of China (Nos. 51822102 and 51731005), and Grant-in-Aid of Scientific Research, JSPS (Kiban S 26220907) is gratefully acknowledged.

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alez, Magnetocaloric effect in preferentially textured Mn50Ni40In10 melt spun ribbons, Appl. Phys. Lett. 94 (2009) 222502, https://doi.org/10.1063/ 1.3147875. [161] J. Liu, N. Scheerbaum, J. Lyubina, O. Gutfleisch, Reversibility of magnetostructural transition and associated magnetocaloric effect in Ni-Mn-In-Co, Appl. Phys. Lett. 93 (2008) 102512, https://doi.org/10.1063/1.2981210. [162] F. Chen, J.L. S
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anchez-Valdés, P. M€ ullner, Y.G. Shi, Y.X. Tong, L. Li, High temperature martensitic transformation and giant magnetocaloric effect in Ni40Co10Mn41Sn9 melt-spun ribbons, J. Alloys Compd. 744 (2018) 493e501, https://doi.org/10.1016/j.jallcom.2018.02.091. [163] Z.Y. Wei, E.K. Liu, J.H. Chen, Y. Li, G.D. Liu, H.Z. Luo, X.K. Xi, H.W. Zhang, W.H. Wang, G.H. Wu, Realization of multifunctional shape-memory ferromagnets in all- d -metal Heusler phases, Appl. Phys. Lett. 107 (2015) 022406, https://doi.org/ 10.1063/1.4927058. [164] Z.Y. Wei, E.K. Liu, Y. Li, X.L. Han, Z.W. Du, H.Z. Luo, G.D. Liu, X.K. Xi, H.W. Zhang, W.H. Wang, G.H. Wu, Magnetostructural martensitic transformations with large volume changes and magneto-strains in all- d -metal Heusler alloys, Appl. Phys. Lett. 109 (2016) 071904, https://doi.org/10.1063/1.4961382. [165] H. Neves Bez, A.K. Pathak, A. Biswas, N. Zarkevich, V. Balema, Y. Mudryk, D.D. Johnson, V.K. Pecharsky, Giant enhancement of the magnetocaloric response in Ni-Co-Mn-Ti by rapid solidification, Acta Mater. 173 (2019) 225e230, https://doi.org/10.1016/j.actamat.2019.05.004.

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[166] M. Kohl, D. Brugger, M. Ohtsuka, T. Takagi, A novel actuation mechanism on the basis of ferromagnetic SMA thin films, Sensors Actuators A Phys 114 (2004) 445e450, https://doi.org/10.1016/J.SNA.2003.11.006. [167] F. Lambrecht, I. Aseguinolaza, V. Chernenko, M. Kohl, Integrated SMA-based NEMS actuator for optical switching, in: 2016 IEEE 29th Int. Conf. Micro Electro Mech. Syst., IEEE, 2016, pp. 79e82, https://doi.org/10.1109/MEMSYS. 2016.7421562. [168] R. Fechner, C. Chlub, E. Quandt, M. Kohl, A shape memory alloy 12 optical waveguide switch, in: 2018 IEEE 18th Int. Conf. Nanotechnol, IEEE, 2018, pp. 1e3, https://doi.org/10.1109/NANO.2018.8626320.

CHAPTER SEVEN

Piezoelectric polymers and composites for multifunctional materials Kuntal Maity1, Dipankar Mandal2 1

Organic Nano-Piezoelectric Device Laboratory, Department of Physics, Jadavpur University, Kolkata, West Bengal, India 2 Institute of Nano Science and Technology, Mohali, Punjab, India

7.1 Introduction The rapid development of modern technology has created significant applications in multiple sectors including smart and wearable electronics (phone, watch, glass, and headset), e-skin sensors, autonomous sensors and actuators, batteries, supercapacitors, physical/chemical sensors, artificial intelligence, optoelectronics, and biomedical systems [1e5]. For example, now a days human health monitoring is not only dependent on hospital-mode; rather, people can monitor real-time health status with wearable electronics. However, the existing technologies suffer from drawbacks such as limitation of power supply (bulky size of the batteries, inconvenience of recharging, and danger of explosions), sensitivity, multifunctional properties, flexibility and wearability, toxicity, and cost-effectiveness [6]. In this context, there is increasing demand in the development of novel and multifunctional materials to provide necessary requirements in modern technology. In addition, the modern technology urgently requires the utilization of environmentfriendly biodegradable materials in the design of accessible and convenient systems as the rapid use of electronics leads to global issues like e-wastes or environmental pollution [7]. Thus it is necessary to develop highly sensitive and functional materials in the integration of modern electronic systems. In this background, polymer materials, particularly being piezoelectric, play a significant role, as they possess versatile mechanical, physical, and chemical properties and can be employed in several applications. Basically the piezoelectric polymers could convert any type of mechanical stress into electricity where the stress can come from a wide range of ambient sources such as wind or acoustic vibrations, any kind of human motions Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00001-9

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(walking, running, stretching, and breathing), noises available everyday [8]. Over the last decade, there is large number of scientific reports about the dielectrics with piezoelectric properties; however, piezoelectric energy harvesting applications have been immensely investigated only after 2006 when Wang et al. put forward a new research field of nanogenerators (NGs) based on zinc oxide (ZnO) nanorod where the requirement of external power supply/batteries can be replaced [9,10]. In addition to the energy conversion ability, the piezoelectric polymers offer versatile applications from energy storing to hydrogen production, filtrations to batteries, thermal to optical sensors, magneto-electric to tissue engineering, particularly in a self-powered manner [11e18]. Furthermore, it is noteworthy to mention that the incorporation of a variety of external fillers such as inorganic piezoelectric materials, for example, lead zirconate titanate (PZT), barium titanate (BaTiO3), zinc oxide (ZnO), natural piezoelectric materials (sugar), carbon-based materials, for example, graphene, carbon nanotubes (CNTs); 2D molybdenum disulfide (MoS2); perovskite methylammonium lead iodide (CH3NH3PbI3); biomaterial deoxyribonucleic acid (DNA) in the polymer matrix have led to interesting functional properties of the polymer composites and thus extends its attractive and promising applications in modern technology [19e26]. Though several investigations on piezoelectric polymers and composites can be found in literature, the collective information on the distinct characteristics and applications in different sectors are uniquely discussed as much as possible in this chapter. A schematic is shown in Fig. 7.1, where the integration of piezoelectric polymers and composites with multifunctional applications is illustrated. The objectives of this chapter are the following: (i) The important properties of the different polymers in the light of piezoelectricity are highlighted along with preparation techniques, advantages, and disadvantages. In particular, polyvinylidene fluoride (PVDF) and its copolymers (for example, poly(vinylidenefluoride-cotri-fluoroethylene) (P(VDF-TrFE)), poly(vinylidene fluoridehexafluoropropylene) (P(VDF-HFP)) exhibiting best all-round activities are presented exclusively. (ii) The state-of-the-art piezoelectric polymer composite with multifunctional ties is covered, where the role of external filler materials are explored. (iii) Particular attention has been paid to recent advancements and growing trends in the modern technology related to multifunctional materials revealing the necessity of current research.

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Bio-medical sectors

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Energy harvesting Figure 7.1 Schematic illustration of piezoelectric polymer and composite for multifunctional applications. Reproduced from K. Maity, S. Garain, K. Henkel, D. Schmeißer, D. Mandal, Self-powered human-health monitoring through aligned PVDF nanofibers interfaced skin-interactive piezoelectric sensor, ACS Appl. Polym. Mater. 2 (2020) 862e878; F.R. Fan, W. Tang, Z. Lin Wang, Flexible nanogenerators for energy harvesting and selfpowered electronics, Adv. Mater. 28 (2016) 4283e4305; K. Maity, D. Mandal, All-organic high-performance piezoelectric nanogenerator with multilayer assembled electrospun nanofiber mats for self-powered multifunctional sensors, ACS Appl. Mater. Interfaces 10 (2018) 18257e18269; K. Maity, S. Garain, K. Henkel, D. Schmeißer, D. Mandal, Natural sugar-assisted, chemically reinforced, highly durable piezoorganic nanogenerator with superior power density for self-powered wearable electronics, ACS Appl. Mater. Interfaces 10 (2018) 44018e44032; S. Garain, S. Jana, T. K. Sinha, D. Mandal, Design of in situ poled Ce3þ-doped electrospun PVDF/Graphene composite nanofibers for fabrication of nanopressure sensor and ultrasensitive acoustic nanogenerator, ACS Appl. Mater. Interfaces 8 (2016) 4532e4540; P. Adhikary, S. Garain, S. Ram, D. Mandal, Flexible hybrid Eu3þ doped P(VDF-HFP) nanocomposite film possess hypersensitive electronic transitions and piezoelectric throughput, J. Polym. Sci., Part B: Polym. Phys. 54 (2016) 2335; C. Ribeiro, S. Moreira, V. Correia, V. Sencadas, J. G. Rocha, F. M. Gama, J. L.G. Ribelles, S. Lanceros-Méndez, Enhanced proliferation of pre-osteoblastic cells by dynamic piezoelectric stimulation, RSC Adv. 2 (2012) 11504e11509; T. Marques-Almeida, V.F. Cardoso, S. Ribeiro, F.M. Gamad, C. Ribeiro, S. Lanceros-Mendez, Tuning myoblast and pre-osteoblast cell adhesion site, orientation and elongation through electroactive micropatterned scaffolds, ACS Appl. Bio Mater. 2 (2019) 1591e1602 with permission from concerning publisher.

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Overall, the piezoelectric polymers and composites as multifunctional materials are described in such approaches which have not been discussed earlier. At last, the requirement of developing multifunctional materials with future research scope is also highlighted in this chapter.

7.2 Piezoelectricity 7.2.1 Piezoelectric phenomenon The term “piezoelectricity” is associated with the Greek word “piezein” which means “to press” or “to squeeze,” i.e., stress-induced electricity generation. The piezoelectricity is typically found in insulators or dielectric materials in which electrical conductivity is very low. The piezoelectric materials exhibit stress-dependent displacement of electric charges or alignment of molecular dipoles which in turn can produce the electricity labeled as the direct piezoelectric effect or experience a mechanical strain under the application of electric field, known as converse piezoelectric effect (Fig. 7.2) [27]. Direct piezoelectric effect Tension Compression

Original shape

0

Converse piezoelectric effect Expansion Contraction

Figure 7.2 Direct piezoelectric effect (electrical signal generation under compression and tension) and converse piezoelectric effect (dimensional expansion and contraction under electrical charge applications) of a piezoelectric material.

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The direct piezoelectric effect is mainly essential for functional materials of nanogenerators, sensors, biomedical applications where the subjected mechanical stresses are utilized to generate electrical charges in the piezoelectric materials whereas converse piezoelectric effect can be found more suitable in actuator, vibration damping, and acoustic emitter applications. Interestingly the piezoelectric materials were first applied during the World War I, but polymers with such properties found practical use much later, in the 1960s to 1970s [28]. It is noteworthy that the properties of piezoelectric materials depend not only on their chemical structure but also on the method and conditions of their manufacturing.

7.2.2 Theoretical bases Piezoelectric materials are associated with the direct and converse piezoelectric effects. These effects are set by the piezoelectric constitutive equations given below [29]:    E   d s s dt ¼ (7.1) T D E d ε where d and s represent the strain and stress components; D and E refer to the electric displacement and electric field components; s, ε, and d are the elastic compliance, the dielectric constant, and the piezoelectric coefficient, respectively; the superscripts E and T denote that the respective constants are evaluated at the constant electric field and constant stress, respectively; and the superscript t stands for the transpose. In general, piezoelectric materials are highly anisotropic, due to their crystal structure. The properties may vary depending on the orientation of the crystal of the materials [30]. The orientation axes (three axes and three rotations) are depicted in Fig. 7.3. Usually axis 3 is placed along the thickness of the sample, and axis 1 is placed in plane along its longest side. All the mechanical and electrical parameters are tensors of different orders. To differentiate the piezoelectric and elastic coefficients measures along different axes, they are represented in matrix form and are symbolized as dij, where index i refers to the direction of electric field and index j refers to the direction of applied stress.

7.2.3 Piezoelectric materials The first discovered piezoelectric material is naturally occurring quartz by French Scientist Jacques and Pierre Curie in 1880 [31]. After about 140 years

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+ –

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Figure 7.3 The coordinate axes employed to define the orientation of a piezoelectric material and illustration of 33 and 31 modes.

of development, many natural piezoelectric materials have been found and synthetic materials are also created with excellent piezoelectric properties for various multifunctional applications (Table 7.1). 7.2.3.1 Inorganic piezoelectric materials Inorganic piezoelectric materials can be classified in two types: piezoelectric crystals and piezoelectric ceramics. Piezoelectric crystals typically have a single crystal structure and exhibit spontaneous piezoelectricity owing to their non-centrosymmetric crystal structure, for example, quartz, Rochelle salt [32]. Whereas piezoelectric ceramics are composed of many small crystals with random crystal orientations and will only show piezoelectricity after a polarization process, normally by applying a high electrical field to align the crystal orientations, such as, BaTiO3, PZT, ZnO, and aluminum nitride (AlN) [33,34]. There are 32 crystallographic classes that depend on the geometry and symmetry of the unit cell, among them 21 are noncentrosymmetric in nature and all but one (due to other symmetry elements) exhibit piezoelectricity. It is the lack of symmetry in the distributions of ions in these crystalline materials that lead to the presence of an electrical dipole exhibiting piezoelectric response. However, the preparation of these single crystals requires a costly precision controlled process, which limits their application in many fields.

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Table 7.1 Different types of typical piezoelectric materials. Types of materials Materials

Natural piezoelectric crystals

Synthetic piezoelectric ceramics Lead-free piezoelectric ceramics Bio-piezoelectric materials

Piezoelectric polymers

Berlinite (AlPO4), sucrose, quartz, Rochelle salt, topaz, tourmaline-group minerals BaTiO3, PZT, AIN, ZnO, PbTiO3, KNbO3, LiNbO3 (K,Na)NbO3, BiFeO3, Bi4Ti3O12, Na0.5Bi0.5TiO3 Bone, tendon, silk, wood, enamel, dentin, DNA, viral proteins PVDF, P(VDF-HFP), P(VDF-TrFE), P(VDF-CTFE), P(VDF-TrFE-CFE), PLLA, odd nylons

7.2.3.2 Bio-piezoelectric materials There are some naturally available materials which show spontaneous piezoelectricity known as bio-piezoelectric materials such as silk, bone, fish scale, specific virus, etc. [35]. Importantly these materials being biocompatible and biodegradable potentially offer a simple and environment-friendly approach to biomedical applications or smart electronics. However, the short lifetime of the bio-piezoelectric materials restricted them from large applications. 7.2.3.3 Piezoelectric polymers In comparison to inorganic piezoelectric materials, polymers materials, such as PVDF and its copolymer P(VDF-TrFE), P(VDF-HFP), polyL-lactide (PLLA) are mechanically flexible, biocompatible, easy for processing, and exhibit promising piezoelectric properties. So these polymer materials are the point of concern for versatile application purposes [36].

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7.3 Piezoelectric polymers 7.3.1 Background Piezoelectric ceramic materials generally show high piezoelectric coefficient and dielectric constant and thus are suitable for high energy harvesting devices or multifunctional applications. But their wide practical applications are restricted due to the brittleness of the materials, complex and cost-intensive processing method, and the presence of toxic materials. On the other hand, bio-piezoelectric materials show spontaneous piezoelectricity, but they suffer from the drawbacks of low piezoelectric coefficient and dielectric constant. In this background, piezoelectric polymers are promising in diverse multifunctional applications because of their superior advantages in mechanical flexibility, easy processing, cost-effectiveness, chemical resistiveness, low density, and biocompatibility over ceramics. Thus piezoelectric polymerbased materials have established various applications in smart and multifunctional systems such as nanogenerators/energy harvesters, sensors, actuators, photodetectors, energy storage devices, and biomedical applications in the form of thin films, nanofibers, foams, textiles, and coatings.

7.3.2 Piezoelectricity in polymers The piezoelectricity in polymers arises due to the high degree of polarization depending on the net molecular dipole moments in the structure. The molecular dipoles in these materials can be polarized or oriented by the application of an external electric field, mechanical stretching, or incorporating external fillers, which in turn leads the piezoelectric response of the polymers. Polymers including vinylidene fluoride (VDF)ebased fluoropolymers, odd-numbered nylon, poly-L-lactide, polyimide show piezoelectricity and have become the choice of materials for multifunctional applications [36e38]. The piezoelectric properties of few polymers are tabulated in Table 7.2. Basically piezoelectric polymers are a class of polymers possessing semicrystalline structure. The piezoelectric performance mainly depends on their dipole orientation in crystalline regions toward the applied poling field, chain conformation, and macromolecular structure. However these polymer materials suffer from the drawbacks of low piezoelectric coefficient and dielectric permittivity than ceramic materials. In this context, different approaches are applied to heighten the piezoelectric properties by modifying

Piezoelectric polymers and composites for multifunctional materials

Table 7.2 Piezoelectric properties of few polymers [36]. Relative permittivity (εr) dij(pCNL1) Polymer

PVDF

6e12

P(VDF-TrFE)

18

P(VDF-HFP)

11

P(VDF-CTFE) 13 P(VDF-TrFE-CFE) 65 Polyamide 11 5 Polyimide ─ 3e4 Poly-L-lactide Polyhydroxybutyrate 2e3.5

d31 ¼ 8 to 22 d33 ¼ 24 to 34 d31 ¼ 12 to 25 d33 ¼ 25 to 40 d31 ¼ 30 d33 ¼ 24 d33 ¼ 140 ─ d33 ¼ 4 d33 ¼ 2.5 to 16.5 d14 ¼ 9.82 d33 ¼ 1.6 to 2.0

247

Electromechanical coupling (k33)

0.20 0.29 0.36 0.36 ─ ─ ─ ─ ─

the molecular structure, forming composites, controlling processing conditions, and post-treatment methods, for instance, mechanical stretching and electrical poling, etc. PVDF and its copolymers are the best eligible polymers for energy harvesting or other multifunctional applications due to their higher piezoelectric coefficient and ease of fabrication than other piezoelectric polymers.

7.3.3 State-of-the-art piezoelectric polymers 7.3.3.1 PVDF Among all the piezoelectric polymers, PVDF has been the mostly investigated polymers for years and found versatile applications in recent years. Structure and Polymorph: PVDF is a semi-crystalline polymer with simple chemical structure of two fluorine atoms at every second carbon atom in the hydrocarbon backbone, i.e., [─CH2─CF2─]n. The polymer retaining about 50% lamellar crystals embedded in amorphous region shows relatively strong piezoelectricity than other polymers. In the crystalline regions, it possesses five distinct crystalline phases such as a, b, g, d, and ε exhibiting a complex structure. Among all the crystalline phases, a, b, and g phases are the most studied and frequently employed polymorphs of PVDF (Fig. 7.4(a)). In more detail, ae phase is electrically nonactive owing to antiparallel arrangement of dipoles (trans-gaucheetrans-gauche, i.e., TGTG/conformation) whereas be (all trans TTTT chain conformations) and ge (T3GT3G/conformations) phases are regarded as electrically active

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(a)

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Hydrogen Fluorine Carbon

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Figure 7.4 (a) Schematic illustration of the chain conformation of the a, b, and g phases for PVDF and (b) Distinct phases identification of the PVDF by XRD and FTIR technique. (c) Schematic presentation of different methods for obtaining the be phase in PVDF. (d) Schematic presentation of ultrathin PVDF films preparation by horizontal LS technique. (e) Highly aligned P(VDF-TrFE) nanofiber preparation by electrospinning technique, corresponding nanofibers are shown in SEM-micrographs. (a,b) reproduced from P. Martins, A.C. Lopes, S. Lanceros-Mendez, Electroactive phases of poly(vinylidene fluoride): determination, processing and applications, Prog. Polym. Sci. 39 (2014) 683e706 with permission by publisher. (d) reproduced from S. Maji, P.K. Sarkar, L. Aggarwal, S.K. Ghosh, D. Mandal, G. Sheet, S. Acharya, Self-oriented bcrystalline phase in the polyvinylidene fluoride ferroelectric and piezo-sensitive ultrathin LangmuireSchaefer film, Phys. Chem. Chem. Phys. 17 (2015) 8159 with permission by publisher. (e) reproduced from L. Persano, C. Dagdeviren, Y. Su, Y. Zhang, S. Girardo, D. Pisignano, Y. Huang, J.A. Rogers, High performance piezoelectric devices based on aligned arrays of nanofibers of poly(vinylidenefluoride-co-trifluoroethylene), Nat. Commun. 4 (2013) 1633e1643 with permission by Publisher.

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in nature as the dipoles of individual molecules are packed parallel to each other leading to a nonzero dipole moment [38]. It is noteworthy to mention that be phase possesses highest polarization showing superior properties like piezo, pyro, and ferroelectric of the materials and thus PVDF is more advantageous than other piezoelectric polymers possessing high piezoelectric coefficient (d33 ¼ 24─34 pC/N; d31 ¼ 8─22 pC/N) [36,37]. Identification of electroactive phases

The existing polymorphs of PVDF can be identified and quantified by different well-established techniques such as Fourier transformed infrared spectroscopy (FT-IR) and X-ray diffraction study. (XRD) and differential scanning calorimetry (DSC) based thermal study. Martins et al. described and summarized the detailed analysis of a, b, and g phases by these techniques [38]. The characteristics of crystalline phases are listed in Table 7.3 for identification. In a brief, a phases can be well identified by FT-IR, but b and g phase vibration bands appear in similar fashion (Fig. 7.4(b)). In contrast, a and g phases exhibit coincident diffraction peaks, but distinct diffraction peak of b-phase is clear in XRD. Thus the crystalline phases can be well identified by the combination of the FT-IR and XRD. Approaches of b- Phase Nucleation: The main problem associated with PVDF is that it is typically stable with nonpolar ae phase. However, it needs to promote into electro-active be and ge phase in order to get multifunctional applications like nanogenerators, sensors and actuators, photodetectors, magneto-electric, energy storage, and tissue engineering applications. Thus many attempts have been made so far to achieve b-phase in PVDF, which we will discuss next (Fig. 7.4(c)). It has been found that b phase can be inducing in PVDF by applying direct stretching mechanism from a-phase. Sencadas et al. reported about Table 7.3 The characteristics of different PVDF polymorphs [38]. Polymorph ae phase be phase

ge phase

FT-IR absorption bands 489,614, 766, 795, 510, 840,1275 431,512, 776,813, 855, 976 833,840,1234 (Wavenumber in cm1) XRD reflections ð2qÞ ( ) 17.6(110), 18.3 20.2(110/200) 18.5 (020), 19.2 and corresponding (020), 19.9(110), (002), 20.1 (110) diffraction planes 26.5(021) 167e172 167e172 179e180,189e190 DSC melting Temperature ( C)

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the transformation from ae to be in PVDF by mechanical stretching at different temperatures [39]. But the quality of the film and piezoelectric phases are compromised so that it is hard to get wide range of applications. Again several researchers have reported about the employing of high pressures more than 4000 kg/cm2 to induce be phase of PVDF from melt [40]. In addition, Scheinbeim et al. also reported that increasing quenching pressures (such as 200e700 MPa) could enhance the remarkable be phase content of the PVDF [41]. Gradys et al. showed the melt-crystallization process of PVDF in nonisothermal mode at ultrahigh cooling rates (30e3000 K/s) at constant temperatures afterward quenching at 6000 K/s [42]. The quenching effects on the crystallization of PVDF from the melt have been investigated in some reports. In another approach, evaporative deposition methods are also interesting for preparing thin films to induce b-phase directly on a substrate particularly suitable for electronic applications. But the evaporation at high temperature (200 C) can decompose molecular chain and further the properties of PVDF [43]. Thus the less-destructive approaches like electrospray technique, ultrasonic atomization, and the use of VDF oligomers instead of PVDF have become popular [44,45]. It is worthy to mention that micrometer order thickness films can be prepared; however the free-stranding films of nanoscale thickness are very hard to fabricate considering previous reports. In this way more attention has been paid to improve more suitable process to get nanoscale films and nanofibers of be PVDF by solvent-casting methods, such as spin-coating, LangmuireBlodgett deposition, solvent evaporation, and electrospinning technique. In this context, spin-coating is a promising process to achieve the be phase of PVDF. Benz et al. reported about the PVDF thin film preparation with high fraction of b-phase by varying spin-speed and the humidity conditions [46]. In another fascinating approach, oriented PVDF ultrathin films can be prepared by LangmuireBlodgett method which helps to obtain bphase. Jiang et al. reported the construction of PVDF ultra-thin films by this technique [47]. In a step further, Mazi et al. reported about the investigation of piezoelectric phase and ferroelectric switching in ultrathin PVDF films adopting horizontal LangmuireSchaefer (LS) method (Fig. 7.4(d)) [48]. They observed that increasing number of LS layers helps to induce b phase exhibiting ferroelectric switching. Besides, employment of high

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electrical field is effective in achieving polar phase in PVDF reported by Davis et al. [49]. However, electrospinning is another facile, scalable, and interesting technique to obtain be phase in PVDF nanofibers (NFs), which is flexible and air-permeable. It should be noted that alignment of molecular dipoles (CH2/CF2) in PVDF chain with b phase induction avoids the employment of further electrical poling [50]. 7.3.3.2 PVDF copolymers Copolymerization is another effective process to tune the polymorph structure as well as the electro-active phase nucleation of PVDF. It is observed that the type of comonomer and composition ratio affect the molecular-structure as well as the ferro-, piezo-, and pyroelectric properties of the polymer. P(VDF-TrFE): P(VDF-TrFE) is a copolymer of PVDF where the introduction of the third fluoride atom in the TrFE [eCHFeCF2] unit helps to induce polar b phase by arranging the molecular chain in all-trans conformation. It has been observed that the nucleation of b phase depends on the amount of TrFE content (typically over 11 mol%) and importantly is independent of the processing conditions or electrical poling (when it is over 20 mol%) [51]. Thus the utilizations of P(VDF-TrFE) in the form of thin films or free standing films or nanofibers have found multifunctional applications by several researchers. Persano et al. reported about the fabrication of flexible nanogenerators and nano-pressure sensors by using electrospun nanofibers of P(VDF-TrFE) (Fig. 7.4(e)) [52]. Importantly, they have shown that the aligned nanofibers could exhibit enhanced polar phase as well as piezoelectric performance. P(VDF-HFP): P(VDF-HFP) is another copolymer of PVDF, where the introduction of bulky -CF3 groups provides more space to allow molecular dipoles to rearrange in polymer chain. The high remnant polarization (w80 mC/m2) and high piezoelectric coefficient (30 pC/N) were found with the HFP content of 5 mol% for solvent casted film [53]. Thus the P(VDF-HFP) shows its promising applications in piezoelectric and ferroelectric areas like the development of magneto-electric sensors and actuators. 7.3.3.3 Other piezoelectric polymers Besides PVDF and its copolymer, there are some other polymers which show the piezoelectric characteristics independently or by incorporating external fillers.

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PLLA: PolyL-lactide (PLLA) is a promising polymer due to its biocompatibility, biodegradability, and fascinating piezoelectric characteristics. Basically PLLA has both direct and inverse piezoelectric effects, but it only exhibits d14 and d25 (i.e., same magnitude and opposite sign) shear piezoelectric constants. PLLA is an optically active polymer consisting of chiral molecules where eCOeO unit is attached to an asymmetric carbon atom and oriented in helical conformation in the polymer chains. The polarization arises from the helix structure of eC═O dipole sheared through its side chain and further shows the piezoelectricity [54]. However the PLLA is stable with randomly oriented C═O dipoles, where chains should be stretched to obtain polar b crystalline phase from nonpolar a crystalline phase. In this context, it is found that electrospinning technique is very effective in the transformation of polarized dipoles (Fig. 7.5(a)). However, PLLA was typically found in biomedical applications by several reports. Zhu et al. described the shear piezoelectricity and showed energy harvesting performance from PLLA nanofibers for smart electronics [55]. In a step further, PLLA nanofiber was also utilized in the fabrication of bio-e-skin sensor fabrication for noninvasive human health signal monitoring reported by Sultana et al. [56]. They observed a longitudinal piezoelectric charge coefficient (d33) of w3  1 pmV1 by piezoelectric force microscopy (PFM), with a height of 650 nm and hysteretic phase switching along with butterflyshaped loops (Fig. 7.5(b)). Polyamides: Polyamides or nylons are a class of polymers possessing semi-crystalline structure, among them odd-numbered nylons exhibit piezoelectricity owing to hydrogen bonding and subsequent dipole arrangement within the polymer chains. In this context, nylon-11 [C11H21ON]n shows distinct polymorph possessing five crystalline structures, including triclinic a and pseudo-hexagonal g phases [57]. The polar crystal structure (g phase) in nylon-11 originates to exhibit piezoelectric characteristics (typically d31 w 3 and 12 pC/N) obtained from melt-drawn, strained, solvent casting, or electrical poling. Kawai first reported about the poled nylon11 films obtained from mechanical stretching and piezoelectric constant (d3l) of w0.5 pC/N [58]. In another process, Datta et al. described about the preparation of self-poled nylon-11 nanowires grown by facile and scalable capillary wetting technique (Fig. 7.6(a)) [59]. Further they used the polar crystalline nanowire (confirmed by XRD and FT-IR, shown in Fig. 7.6(b)) in the fabrication of nanogenerator producing electrical voltage of 1 V and short-circuit current of 100 nA under the application of external mechanical pressure.

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Figure 7.5 (a) Molecular orientation of PLLA chain before and after electrospinning technique. (b) (i) AFM topography and (ii) amplitudee voltage hysteresis loops of the PLLA nanofibers. Reproduced from A. Sultana, S.K. Ghosh, V. Sencadas, T. Zheng, M.J. Higgins, T.R. Middya, D. Mandal, Human skin interactive self-powered wearable piezoelectric bio-e-skin by electrospun poly-L-lactic acid nanofibers for non-invasive physiological signal monitoring, J. Mater. Chem. B 5 (2017) 7352e7359 with permission from publisher. 253

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Figure 7.6 (a) Schematic presentation of the preparation of nylon-11 nanowires via AAO template. (b) XRD and FT-IR spectra of nylon-11 nanowires. Reproduced from A. Datta, Y.S. Choi, E. Chalmers, C. Ou, S. Kar-Narayan, Piezoelectric nylon-11 nanowire arrays grown by template wetting for vibrational energy harvesting applications, Adv. Funct. Mater. 27 (2017) 1604262 with permission from publisher.

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Besides, other polymers such as polyimide (PI) obtaining polar functional groups also show piezoelectricity and have found in sensor applications owing to its high thermal resistance (w400 C) [60].

7.4 Piezoelectric polymer composites Till now we have described about different methods in nucleating polar phase in piezoelectric polymer; however, there exist some drawbacks such as, structural deformations or microstructural limitations constraining some applications like electro-optical sensors and nonvolatile memories [61,62]. Thereafter, researchers have found alternative approach by incorporating different fillers such as ceramic materials (BaTiO3, PZT), ionic liquids, carbon materials (carbon nanotube, graphene), metal nanoparticles (Pd, Ag, Pt), metal salts, clay, hydrated ionic salts, PMMA [63e69]. It is noteworthy to mention that the incorporation of fillers leads to induce/enhance the polar content in polymer materials by the nucleation mechanism and interestingly the polymer composite shows distinct properties depending on the functional properties of fillers. Thus the polymer composite materials are point of attraction having multifunctional applications which we discuss in the next section.

7.4.1 PVDF composites In this background, PVDF composites are the mostly studied and employed materials for multifunctional applications, particularly for energy harvesting and energy storage purposes. Starting with the ceramic fillers used in the preparation of PVDF nanocomposites, Mendes et al. described about the polar b-phase nucleation in PVDF influenced by distinct sizes of BaTiO3 particle [63]. Again, it has been found that the incorporation of clay into PVDF matrix could induce polar phases in PVDF. Patro et al. reported about the preparation of PVDF-clay nanocomposites via one melt-mixing process and observed that highest polar phase (w99% b phase) nucleation by doping with phosphonium surfactant modified clay [70]. After that PVDF-montmorillonite, PVDF-graphite nanosheets, PVDF-palladium nanoparticle (Pd NP), PVDF-gold nanoparticle (Au NP) composites are reported becoming the alternative approach to induce polar phases in PVDF [71e74]. Mandal et al. reported about the Pd NP-doped PVDF thin film where b phase has been promoted due to the interaction between negative surface charge of Pd NPs and positive end of PVDF chains [73]. The respective FE-SEM images exhibit dominant a spherulitic growth in

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neat PVDF whereas b crystallites are more dominant in Pd NPs doped PVDF composites (Fig. 7.7(a)). In a similar approach, PVDF composite films are prepared by doping Au NPs where it has been shown that the doping of higher concentration of Au NPs promotes the b phase as well as crystalline polymorph as evident from the respective FT-IR spectra and topographical image (Fig. 7.7(b)) [74]. Besides, there are reports about the employment of rare-earth hydrated salts in the PVDF where the existence of water helps to induce polar phase by forming hydrogen bonding with polar groups in PVDF. In this background, Garain et al. reported about the high content of b phase (99%) into cerium(III)eN,N-dimethylformamide-bisulfate [Ce(DMF) (HSO4)3] complex-PVDF composite films and studied the self-poled mechanism behind the polar phase nucleation (Fig. 7.7(c)) [68]. They observed that the composite exhibits UV-light emitting due to the existence of Ce3þ ions in the composite and further utilized in the energy harvesting applications. In a step further, several researchers reported about the fabrication of PVDF composite nanofibers by incorporating external fillers for multifunctional use. Here the role of fillers is crucial as it stabilizes the b phase formed by electrospinning technique and further enhances the polar phase content in polymer composites. In this regard, Maity et al. [24] for the first time reported about the fabrication of PVDF nanofibers by incorporating 2D transition metal dichalcogenide (TMD) MoS2 which is prepared by simple exfoliation method (shown in the schematic of Fig. 7.8(a)). They have observed that doping of external fillers in the PVDF matrix leads to enhance the b phase formation. The corresponding surface morphological structure of PVDF/MoS2 nanofibers and MoS2 lattice fringes realized by high-resolution transmission electron microscopy (HR-TEM) is shown in Fig. 7.8(b). In a similar approach, natural piezoelectric material sugar is encapsulated by PVDF matrix reported by Maity et al. [21]. They described that incorporation of sugar into PVDF leads to reduce the diameter of nanofibers as clear from the surface structure and respective histogram profiles (Fig. 7.8(c)) and enhanced the polar content. Thus it was expected that the composite nanofibers could show promising piezoelectric performance. Liu et al. reported about the graphene oxide (GO) wrapped PVDF nanofibers, where the significant enhancement of piezo-response was realized by hysteresis loops and polarization reversal in comparison to neat PVDF NFs (Fig. 7.8(d)) [75].

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Figure 7.7 (a) FE-SEM images of neat PVDF and Pd-NP-doped PVDF thin films. (b) FT-IR spectra of neat PVDF and PVDF-Au composite films. (c) Schematic presentation of cerium complex (CEPLX)-PVDF composite interaction and mechanism. (a) reproduced from D. Mandal, K.J. Kim, J.S. Lee, Simple synthesis of palladium nanoparticles, b-phase formation, and the control of chain and dipole orientations in palladium-doped poly(vinylidene-fluoride) thin films, Langmuir 28 (2012) 10310e10217 with permission by publisher. (b) reproduced from D. Mandal, K. Henkel, D. Schmeißer, The electroactive b-phase formation in Poly(vinylidene fluoride) by gold nanoparticles doping, Mater. Lett. 73 (2012) 123e125. with permission by publisher. (c) reproduced from S. Garain, T.K. Sinha, P. Adhikary, K. Henkel, S. Sen, S. Ram, C. Sinha, D. Schmeißer, D. Mandal, Selfpoled transparent and flexible UV light-emitting cerium complexePVDF composite: a high-performance nanogenerator, ACS Appl. Mater. Interfaces 7 (2015) 1298e1307 with permission from publisher.

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Figure 7.8 (a) Schematic illustration of MoS2 nanosheets dispersed in DMF and the preparation of the PVDF-MoS2 composite nanofiber. (b) TEM image of MPNFW and HRTEM image exhibits the MoS2 fringes present on the PVDF nanofiber surface. (c) FE-SEM images of PNFW and SGNFW; the diameter distributions of the NFWs are shown in the respective inset; scale bars correspond to 5 mm. (d) PFM phase and amplitude responses for the PVDF/GO and pure PVDF nanofibers. (a,b) reproduced from K. Maity, B. Mahanty, T.K. Sinha, S. Garain, A. Biswas, S. K. Ghosh, S. Manna, S.K. Ray, D. Mandal, Two-Dimensional piezoelectric MoS2-modulated nanogenerator and nanosensor made of poly(vinlydine fluoride) nanofiber webs for self-powered electronics and robotics, Energy Technol. 5 (2017) 234e243 with permission from publisher. (c) reproduced from K. Maity, S. Garain, K. Henkel, D. Schmeißer, D. Mandal, Natural sugar-assisted, chemically reinforced, highly durable piezoorganic nanogenerator with superior power density for self-powered wearable electronics, ACS Appl. Mater. Interfaces 10 (2018) 4401844032 with permission by publisher. (d) reproduced from X. Liu, J. Ma, X. Wu, L. Lin, X. Wang, Polymeric nanofibers with ultrahigh piezoelectricity via self-orientation of nanocrystals, ACS Nano 11 (2017) 1901e1910 with permission from publisher.

7.4.2 PVDF copolymers and composites In recent years, PVDF copolymer composites have also become of research interest similar like PVDF composites. Different fillers like SWNT, carbon black (CB) were introduced in the P(VDF-HFP) composite and investigated the piezoelectric behavior of the respective composites [76,77]. Batth et al. found highest piezo-response at a particular loading of SWNT whereas the P(VDF-HFP)/CB composite film was found to be enhanced piezoelectric

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output response at a particular concentration reported by Wu et al. [76,77]. Again, Adhikary et al. reported about the spongelike microporous electroactive P(VDF HFP) film by the adding of hydrated salt MgCl2, 6H2O [78]. In a step further these authors reported about the rare earth ion Eu3þ doped P(VDF-HFP) hybrid film suitable for light emission and mechanical energy harvesting purposes [79]. They showed that the incorporation of Eu3þ induces the polar phase in PVDF-HFP as clear from the crystallographic investigation in Fig. 7.9(a) (i) and the corresponding FE-SEM images revealed that the a spherulites were diminished in the composite film transforming from neat film shown in Fig. 7.9(a) (ii and iii). Another well-known copolymer of PVDF, i.e., P(VDF-TrFE) is also adequately studied and further employed in the preparation of composites. In 2012, Dodds et al. [80] reported about the P(VDF-TrFE) thin film preparation by adding ZnO-NPs and studied the increased remnant polarization (5.8e15.2 m/cm2 due to of 60 MV/m electric field application) by increasing the concentration of ZnO-NPs. Chuan et al. [81] investigated the dielectric properties of P(VDF-TrFE) composite by incorporating ceramic filler like PZT. Further, Batra et al. reported about the multiwalled carbon nanotube (MWCNT)-doped P(VDF-TrFE)/PZT film and studied the enhanced dielectric behavior [82]. In a step further, Ghosh et al. reported about the P(VDF-TrFE) composite film introducing magnetostrictive nickel ferrite (NiFe2O4) NPs [83]. They showed the existence of NPs with lattice plane arising in the composite film containing higher content of NPs as clear from XRD spectra (Fig. 7.9(b) (i)). In addition, the enhancement of b phase with higher loading of NPs in P(VDF-TrFE) was clear from FT-IR spectra (Fig. 7.9(b) (ii)). As a result the composite exhibited enhanced piezoelectric response and thus magneto-electric performance suitable for sensor applications.

7.4.3 Other polymer composites It is noteworthy that besides PVDF and PVDF copolymer composites, other polymer composites possessing piezoelectric characteristics have been studied by several researchers. Basically, piezoelectric ceramic materials (such as PZT, BaTiO3, etc.), metal/semiconductor NPs (Ag, Au, ZnS, etc.), natural crystals are incorporated into the polymer matrix, and interestingly, the composite exhibits multifunctional properties and compatibility. In addition, these polymer composites are alternative to inorganic, brittle piezoelectric materials which are avoided in practical cases, particularly in biomedical applications.

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Sakamoto et al. [84] reported about the incorporation of ceramic filler PZT into polyurethane (PU) and the composite of PU/PZT showed piezoelectric value (d33)  24 pC/N and promised as functional material in sensor and actuator applications. David et al. [85] reported about the preparation polyamide 11 (PA 11)/sodium niobate nanowire (NW) composites film. The composite film exhibited piezoelectric property and would be suitable for piezoelectric sensor applications. Khanbareh et al. [86] added lead titanate (PT) ceramic materials into polyethylene oxide (PEO) and prepared a sensitive and flexible composite. Interestingly the PT-PEO composite exhibited enhanced electrical conductivity, as well as piezoelectric and pyroelectric properties. There is report by sultana et al. about the preparation of electrospun nanofiber composite by the incorporation of piezoelectric semiconductor, i.e., oriented zinc sulfide (ZnS) nanorods (NRs) in semicrystalline poly(vinyl alcohol) (PVA) matrix [87]. The composite shows piezoelectric property due to the existence of noncentrosymmetric ZnS and further used in the mechanical energy harvesting applications being large area mat for polymer materials. Further Zhao et al. [88] prepared a carbonized polyacrylonitrile (PAN-C) and barium titanate (BTO) composite nanofiber mat for self-powered flexible sensor capable of detecting pressure and curvature by integrating piezoresistive, piezoelectric, and triboelectric effects.

7.5 Multifunctional applications In the previous sections, it is found that subsequent development has been made in the processing of piezoelectric functional materials. Thus piezoelectric polymers and composite materials have tremendous applications in multifunctional sectors, particularly in energy harvesting and storage technology, sensors and actuators, photodetectors, optical sensors, and biomedical sectors owing to their outstanding properties, which we will discuss briefly in the following sections.

7.5.1 Mechanical energy harvesters The polymer-based piezoelectric materials are integrated into devices and aimed at converting mechanical energy (e.g., wind, friction, acoustic, ultrasonic waves, and the flow of fluids) into electrical energy, owing to their electromechanical conversion properties. The generated electricity can be employed in powering the portable, smart electronic devices where the power requirement is typically in mW/nW range. Noteworthy to mention that

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the harvesting energy directly from ambient sources is one of the most promising approaches in order to power nanoscale devices without the need of external power supplies. Here we will highlight some recent advancement of NGs for mechanical energy harvesting for building selfpowered systems. The first PVDF NF-based NG was fabricated using metallic electrodes reported by Chang et al. After that, Fang et al. also prepared NG utilizing nonwoven PVDF NFs by aluminum (Al)-based electrodes. Later Dhakras et al. also prepared device by attaching copper (Cu) electrode [11]. Subsequently, Maity et al. [21] reported about the natural sugar-PVDF composite NFs-based organic NG recently. Interestingly the NG exhibited high output voltage and current (100 V, 6.2 mA under 10 kPa) under the external stress applications by human finger touch, drum beats in comparison to neat PVDF NFs-based NG (Fig. 7.10(a) (i, ii, iii and iv)). They also observed that the NG was able to produce high output power density for self-powered electronics devices such as light emitting diodes (LEDs), liquid crystal display (LCD) screen (Fig. 7.10(a) (v and vi)). Furthermore, they demonstrated about the wind energy harvesting by this NG (a schematic of the experimental setup shown in Fig. 7.10(b) (i)). The output voltage responses are shown in Fig. 7.10(b) (ii, iii and iv) obtained by varying experimental parameters such as different distances, air velocities, and effective areas of the NG, respectively. The results promise the utilization of wind vibration which is abundantly available in our environment in order to generate electricity. Dhakras et al. reported about the organiceinorganic PNG by the combination of P(VDF-TrFE) NFs and BaTiO3 NF paste [89]. Further they demonstrated that the high output power was employed in the charging of a mobile phone (Fig. 7.10(c)). Adhikary et al. also reported about the PNG based on Zn2þ/PVDF-HFP composite film for mechanical energy harvesting application with UV-light sensing in self-power mode [90].

7.5.2 Sensors and actuators In recent years the developments of flexible and wearable sensors and actuators have attracted significant interest. In this context, it has been found that PNGs with high sensitivity are very much suitable for the design of selfpowered sensor applications as the existing sensors mainly depend on the external power supply. Different fillers are utilized in the piezoelectric polymer to enhance the piezoelectric property as well as the sensitivity. Thus typically piezoelectric polymer composites exhibit distinct functional properties suitable for sensor designing.

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Figure 7.10 (a) Output voltage responses with different applied human finger impacts for the (i) sugarePVDF NF-based NG, (ii) Ref. NG, (iii) the impact of drum beats of a toy, and (iv) measured short-circuit output current response under the human finger impact of 10 kPa; (v) dependences of output voltage and current on variable external load resistance, and (vi) instantaneous output power density of the NG where the insets showing the lit up array of several LEDs by direct finger touch without using any external power supply. (b) Wind energy harvesting: (i) schematic illustration of the wind blowing setup from an air blower and continuous data acquisition from a computer control system. The real-time voltage signal response of the NG to wind blowing with (ii) different distances, (iii) different air velocities, and (iv) different effective areas of the NG (the corresponding insets show the voltages as a function of distance, velocity, and effective area, respectively). (c) Charging of a mobile by NG; the left image reveals the powering of the device with a vibrator and the right image reveals that mobile charging could be started. (a,b) reproduced from K. Maity, S. Garain, K. Henkel, D. Schmeißer, D. Mandal, Natural sugar-assisted, chemically reinforced, highly durable piezoorganic nanogenerator with superior power density for self-powered wearable electronics, ACS Appl. Mater. Interfaces 10 (2018) 44018e44032. with permission by publisher. (b) reproduced from D. Dhakras, S. Ogale, High-performance organiceinorganic hybrid piezo-nanogenerator via interface enhanced polarization effects for self-powered electronic systems, Adv. Mater. Interfaces (2016) 1600492 with permission from publisher.

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Garain et al. described about the acoustic sensor fabrication based on Ce3þ doped PVDF/graphene composite NFs along with mechanical energy harvesting capability. They demonstrated that the NG could generate acoustic energy under the sound pressure level (SPL) of 88 dB and lit up several LEDs (Fig. 7.11(a) (i)) [22]. Subsequently they stored this energy into commercially available different capacitors, i.e., 1,2.2, and 4.7 mF (Fig. 7.11(a) (ii)). Further they showed that the NG is sensitive to different musical vibrations like piano, guitar, and violin (Fig. 7.11(a) (iii)). The NG also exhibited frequency sensitivity when voltage was increased at a particular frequency (from 110 to 130 Hz) shown in Fig. 7.11(a) (iv). In a step further, Maity et al. [24] reported about the speech signal monitoring system by acoustic sensor based on MoS2/PVDF NFs mat where distinct wave responses for pronouncing of different English letters (e.g., A, B, C, D, and E) from a speaker were subtly captured by this sensor (Fig. 7.11(b) (i)). They further analyzed the spectrum by fast Fourier transform (FFT) for each wave-form and found distinct amplitude peaks (fp) at different frequencies, such as 139.6, 133.4, 142.6, 130.4, and 136.6 Hz, during the pronunciation of A, B, C, D, and E, respectively (Fig. 7.11(b) (ii)). The time-dependent FFT spectra (Fig. 7.11(b) (iii)) also confirmed the sensing performance of different acoustic responses obtained from different letters and thus the sensor was expected to show promising applications in security sectors. In another report by Maity et al. [11] a tactile sensor was fabricated by layer-by-layer structure of PVDF NFs mat and utilized in weight measurement and vibration capturing purposes. For these purposes, they demonstrated that different persons with distinct weights (47.8, 56, 60.6, 62.3, 71.4, and 80 kg) walked upon this sensor and captured the responses (Fig. 7.12(a)). Evidently, the sensors could distinguish the responses and further analyzed by principal component analysis (PCA) (Fig. 7.12(a)). After that they also showed that different vibrations from different household devices (such as mobile, sewing machine, kitchen blender) were capable to detect as well as generate energy (Fig. 7.12(b)). In another approach, Maity et al. [5] reported about the skin-interactive piezoelectric sensor based on aligned PVDF NFs mat where human health monitoring was realized in self-powered manner. Due to the ultrasensitivity, flexibility, and highconformability, the sensor was attached upon the human wrist, neck, arm and throat and subsequent responses were captured during bending and compression (Fig. 7.13(a) (i, ii, iii, and iv)). The coughing actions and throat movement responses during drinking were distinctly captured by this sensor

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Figure 7.11 (a) (i) Output response (illumination of 3 LEDs shown in the inset), (ii) capacitor charging response of different capacitors driven from the music (pop song), (iii) responses with various instrumental (flute, guitar, and violin) and (iv) increasing frequency response and frequency driven 3 LEDs lit up shown in the inset. (b) (i) Output responses observed from the PNG as different letters (A, B, C, D, and E) are pronounced, (ii) the corresponding FFT-processed frequency spectrum and (iii) FFT-processed time-dependent spectrogram (the left side scale bar denotes the amplitude). (a) reproduced from S. Garain, S. Jana, T.K. Sinha, D. Mandal, Design of in situ poled Ce3þ-doped electrospun PVDF/Graphene composite nanofibers for fabrication of nanopressure sensor and ultrasensitive acoustic nanogenerator, ACS Appl. Mater. Interfaces 8 (2016) 4532e4540 with permission by publisher. (b) reproduced from K. Maity, B. Mahanty, T.K. Sinha, S. Garain, A. Biswas, S.K. Ghosh, S. Manna, S.K. Ray, D. Mandal, Two-Dimensional piezoelectric MoS2-modulated nanogenerator and nanosensor made of poly(vinlydine fluoride) nanofiber webs for self-powered electronics and robotics, Energy Technol. 5 (2017) 234e243 with permission from publisher.

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Figure 7.12 (a) Schematic of the weight measurement mapping by walking of different people (different weights) for 10 times individually upon tactile sensor for the purpose of studying the capturing response and respective PCA 2D plot. (b) Vibration sensor with different household device (i) output signal responses of a mobile vibration (during off/on), (ii) machine vibration responses employing sewing machine and (iii) kitchen blender vibration responses. Reproduced from K. Maity, D. Mandal, All-organic highperformance piezoelectric nanogenerator with multilayer assembled electrospun nanofiber mats for self-powered multifunctional sensors, ACS Appl. Mater. Interfaces 10 (2018) 18257e18269 with permission from publisher.

(Fig. 7.13(a) (v and vi)). Further they demonstrated the vocal cord vibration responses during the pronunciation of word “NANO” and different words “N,” “A,” and “O” (Fig. 7.13(b) (i and ii)). The output responses were further analyzed by short time FFT (STFFT) (Fig. 7.13(b) (iii)) and distinguished so that the sensor promised significant applications in the diagnosis of a patient’s damaged vocal cords, throat infection, or any abnormal disorder. In a step further, the authors showed that the sensor could easily capture human wrist pulse signal where a-wave, b-wave, c-wave, d-wave are four systolic waves and e-wave is one diastolic wave (Fig. 7.13(b) (iv and v)) and the respective STFT spectra (Fig. 7.13(b) (vi)) revealed the frequency range of 0e20 Hz of a heart pulse for a healthy person.

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Figure 7.13 Self-powered human health status monitoring by e-skin sensor: (a) real-time sensor responses during (i) wrist bending, (ii) neck and (iii) arm movements, (iv) drinking water, (v) coughing, and (vi) swallowing conditions with enlarged views of the marked area and photographs in the insets, respectively; (b) (i) repeated speaking of the word “Nano,” (ii) speaking of individual letters “N”, “A,” “N,” “O,” (iii) STFT-processed 3D spectrogram of the response related to the speaking of the word “Nano,” (iv) radial artery pulse measurement response by simply interfacing the e-skin sensor to the human wrist with the enlarged view of one marked cycle shown in (v), and (vi) respective STFT-processed 3D spectrograms. Reproduced from K. Maity, S. Garain, K. Henkel, D. Schmeißer, D. Mandal, Selfpowered human-health monitoring through aligned PVDF nanofibers interfaced skin-interactive piezoelectric sensor, ACS Appl. Polym. Mater. 2 (2020) 862e878 with permission from publisher.

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Pérez et al. described about the PVDF-based bimorph actuator where high speed laser beam manipulation was possible owing to the structural and electromechanical properties of the PVDF and further suitable for laser scanner/switcher [91].

7.5.3 Optical sensor/photodetector In the era of smart-sensor systems, optical sensors or photodetectors are widely employed in multifunctional applications such as infrared thermal imaging and remote sensing, ray-measurement/detection, industrial automatic control, photometric measurement, missile guidance, etc. [92]. In this background, there is an increasing demand for the development of self-powered optical sensor or photodetector where the requirement of batteries could be avoided. Thus piezoelectric polymer composites are simultaneously developed and integrated into such sensor systems. Adhikary et al. fabricated novel red light-emitter composed of Eu3þ/ PVDF-HFP film where the incorporation of Eu3þ ions in the polymer made the composite possessing of hypersensitive electronic transitions [79]. The visible light emission spectra observed for the composite film originating due to the transition of 5D0 e 7Fj0 where j0 ¼ 0─4 (Fig. 7.14(a) (i and ii)). Further they evidenced the red light-emission confirmed by fluorescent microscopy image and the Commission Internationale de l’Eclairage (CIE) 1931 chart opening the potential in optical switching/display technology (Fig. 7.14(a) (iii and iv)). In a step further this authors fabricated a PVDF-HFP/Zn2þ composite film based self-powered UV light sensor (Fig. 7.14(b)) [90]. The change of current with fixed bias voltage under UV light on/off situation was confirmed by the currentevoltage (IeV) characteristics. Further the UV light sensing performance in self-powered mode was observed by decrement of current (from 6 to 3.5 nA) under UV light illumination and continuous stress application (Fig. 7.14(b)). In another report Sinha et al. reported about the self-powered optical sensor composed of graphene-silver-PVDF film [93]. The plasmonic characteristics of the composite film in self-powered manner were shown by change in piezo-voltage (w50%) and piezo-current (w70%) under the visible light exposure and mechanical stress applications due to the presence of silver (Ag) NPs (Fig. 7.15(a)). Thus the composite film paves significant potential in futuristic self-powered optoelectronic-sensors. In a step further a self-powered photodetector was fabricated based on methylammonium lead iodide (CH3NH3PbI3) (MAPI)/PVDF composite films reported by

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Figure 7.14 (a) (i) Light emission spectra, (ii) energy level diagram respective of the light emission, (iii) fluorescence microscopy image, and (iv) CIE diagram as per spectrum of the Eu3þ:c-HFP films. (b) Schematic presentation for UV sensor application made with Zn2þeHFP film with the IeV characteristics under UV light illumination ON /OFF and the change of the piezo-current under UV light OFF/ON conditions with the continuous 14 kPa of stress applications. (a) reproduced from P. Adhikary, S. Garain, S. Ram, D. Mandal, Flexible hybrid Eu3þ doped P(VDFHFP) nanocomposite film possess hypersensitive electronic transitions and piezoelectric throughput, J. Polym. Sci., Part B Polym. Phys. 54 (2016) 2335 with permission by publisher. (b) reproduced from P. Adhikary, D. Mandal, Enhanced electro-active phase in a luminescent P(VDFeHFP)/ Zn2þ flexible composite film for piezoelectric based energy harvesting applications and self-powered UV light detection, Phys. Chem. Chem. Phys. 19 (2017) 17789e17798 with permission by publisher.

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Figure 7.15 (a) (i) Change of piezo-response in terms of output voltage, (ii) capacitor charging performance, (iii) change in output current, and (iv) output power of plasmonic sensor with variable resistance. (b) Photodetector operations: (i) currentevoltage (IeV) characteristics under dark and different light illuminations, (ii) current versus time curves with multiple cycle operation and an enlarged view of photocurrent for rise and decay time calculation is shown in (iv) and (iii) photocurrent response under illumination of light of different wavelengths, (iv) The rising and decay time of photoresponse is illustrated. (a) reproduced from T.K. Sinha, S.K. Ghosh, R. Maiti, S. Jana, B. Adhikari, D. Mandal, S.K. Ray, Graphene-silver-induced self-polarized PVDF-based flexible plasmonic nanogenerator toward the realization for new class of selfpowered optical sensor, ACS Appl. Mater. Interfaces 8 (2016) 14986e14993 with permission by publisher. (b) Reproduced from A. Sultana, P. Sadhukhan, M.M. Alam, S. Das, T.R. Middya, D. Mandal, Organo-Lead halide perovskite induced electroactive b-Phase in porous PVDF films: an excellent material for photoactive piezoelectric energy harvester and photodetector, ACS Appl. Mater. Interfaces 10 (2018) 4121e4130 with permission from publisher.

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Sultana et al. [25]. They observed the fast rising time and instant increase in the current under light illumination owing to the presence of perovskite MAPI in the composite film indicating significant potential in the photo detection sector (Fig. 7.15(b)).

7.5.4 Biomedical applications In the present era, the developments of novel materials with biological functionalities are in high demand for its applications in biomedical sectors. Natural materials such as collagen or chitosan are widely used, but they are restricted in specific applications owing to their nonversatile mechanical, physical, and chemical properties. In this regard, biocompatible polymers are mostly studied materials for the employment as biomaterials in tissue engineering, biosensors, smart prostheses purposes. In more recent, smart, and functional materials particularly piezoelectric polymers are well recognized having capability to convert thermal, mechanical and magnetic fields into electrical response [94]. Thus these materials are promising to be employed as smart scaffolds to stimulate cell growth, biosensors, and biocompatible materials. Ribeiro et al. investigated the biological response of cells for poled PVDF (b-PVDF) films [95]. They studied the MC3T3-E1 osteoblast cell culture and noted that the b-PVDF film exhibited higher performance (osteoblast adhesion and proliferation w500 cell/mm2), owing to the existence of surface charge under dynamic conditions than nonpoled samples. The respective cell-culture results for different PVDF films captured by LIVE/DEAD assay exhibited no dead cells after 3 days (Fig. 7.16(a)). Interestingly, this study indicated that piezoelectric PVDF films were effective for cell growth and proliferation due to supply of electrical stimuli and extended applications in composites. Thereafter several reports of polymer composites have come into notice utilizing external fillers such as Ag NPs in antimicrobial activities, zeolites in drug-release, ferrite NPs, and multiferroic composites for cell stimulation [96e99]. Further Damaraju et al. utilized electrospun PVDF NFs as scaffolds for bone tissue engineering [100]. They prepared NFs at different applied electrical voltage (12─30 kV) owing to the variation of b phase content in PVDF by electrospinning technique and utilized NFs prepared at 12 and 25 kV (PVDF-12 kV and PVDF-25 kV correspondingly) for the osteogenic differentiation of human mesenchymal stem cells (MSCs). The confocal microscopy images obtained to investigate the morphologies of MSCs for respective scaffolds exhibited that the cells were attached and well spread

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on both scaffolds (shown in Fig. 7.16(b)). Interestingly they observed that cells on PVDF-25 kV scaffolds had the highest alkaline-phosphatase activity and early mineralization between 10 days as compared to PVDF-12 kV and other polymer materials. In a step further Marques-Almeida et al. reported about the investigation of micropatterned P(VDF-TrFE) scaffolds obtained by soft-lithography technique having distinct surface morphologies and topographies (lines, intermittent lines, hexagons, linear zigzags, and curved zigzags with 25, 75, and 150 mm dimensions, respectively) for the employment in two different cell lines, i.e., myoblast and preosteoblasts [101]. Noteworthy to mention that electroactive P(VDF-TrFE) was used due to its attractive mechanical, physical and chemical properties (spontaneous b phase). They studied the cell-adhesion assays by fluorescent images of C2C12 myoblast cells and preosteoblasts MC3T3-E1 cells on scaffolds (dense and porous) with two topographies of hexagons and lines (Fig. 7.17(a and b)). They observed that the significant improvement in the regeneration of musculoskeletal tissue was exhibited by linear surface topographies and dense morphologies whereas the maximum growth and regeneration of bone tissue was found by the utilization of nonpatterned or anisotropic surface structure-based scaffolds. The promising results of piezoelectric polymers lead to extend the future research in this biomedical sector including sterilization and cleaning procedures, biocompatibility, and development of multifunctional polymer materials, particularly applicable in living organisms. Thus the idea will continue to enhance the scientific innovations in the field of piezoelectric polymer-based biomedical systems.

7.6 Conclusion and perspectives In summary, different polymers (particularly PVDF and its copolymers) and their composites with piezoelectric characteristics are discussed in the light of multifunctional applications. In recent years, the piezoelectric polymers and composites are finding versatile applications including energy harvesting and storage technologies, sensors and actuators, photodetectors and biomedical sectors rapidly owing to their variety of advantages which are not found in other materials. Most importantly the piezoelectric polymers stand out as an environment friendly and cost-effective alternative to inorganic materials possessing toxic elements, such as lead. The piezoelectric efficiency along with functional property can be effectively tuned by

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adopting distinct approaches including proper preparation techniques (such as, nanofiber fabrication by electrospinning, free standing films by solvent casting), structure modifications, (copolymerization) or introducing different fillers. Therefore it becomes very interesting to study the fundamental principle of piezoelectric effect in polymers, as well as the composites. The multifunctional applications provide a new platform to explore the polymer composite with novel materials such as 2D materials or perovskite materials possessing a wide range of unique electrical, optical, mechanical, and thermal properties. Varieties of multifunctional applications have been demonstrated so far by using piezoelectricity in polymers and composites. However there are plenty of scopes to be explored for future fundamental research or industrial applications, such as (i) There are materials that are still not investigated in the polymer composite form such as 2D materials to perovskite materials or biomaterials (very limited have been reported so far). Therefore it is required to study the influence of such materials in the electronic structure of piezoelectric polymers as well as the functional property. (ii) Polymer materials with biological functionalities are in high demand from filtration to cell growth and proliferation or antimicrobial activities. For example, amidst the novel corona virus outbreak, it is highly needed to fabricate masks with efficient air-filter materials which can inhibit the small virus particles (typically in mm range) with high airfiltration efficiency or detoxify the viruses with biofunctionality. (iii) It is needed to optimize the material properties by changing fillers in terms of content, shape and size, copolymer structures (for example, P(VDF-TrFE), P(VDF-HFP)) and so on the multifunctional properties, particularly in industrial-scale which is not explored very much till now. Furthermore, there are also piezoelectric polymers other than PVDF and copolymers (for example, odd nylons, PLAA, polyuria, poly(vinyl fluoride), poly(vinyl chloride), etc.) and thus their composites with multifunctionalities can be studied with much attention in future. At last, it can be concluded that piezoelectric polymers and composites are still at the base of impactful, dynamic, and interesting research field, which will surely bring more innovations in challenging areas in near future.

Acknowledgments We acknowledge the financial support of grant (EEQ/2018/001130) from the Science and Engineering Research Board (SERB), Government of India.

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[21] K. Maity, S. Garain, K. Henkel, D. Schmeißer, D. Mandal, Natural sugar-assisted, chemically reinforced, highly durable piezoorganic nanogenerator with superior power density for self-powered wearable electronics, ACS Appl. Mater. Interfaces 10 (2018) 44018e44032. [22] S. Garain, S. Jana, T.K. Sinha, D. Mandal, Design of in situ poled Ce3þ-doped electrospun PVDF/Graphene composite nanofibers for fabrication of nanopressure sensor and ultrasensitive acoustic nanogenerator, ACS Appl. Mater. Interfaces 8 (2016) 4532e4540. [23] S. Huang, W.A. Yee, W.C. Tjiu, Y. Liu, M. Kotaki, Y.C.F. Boey, J. Ma, T. Liu, X. Lu, Electrospinning of polyvinylidene difluoride with carbon nanotubes: synergistic effects of extensional force and interfacial interaction on crystalline structures, Langmuir 24 (2008) 13621e13626. [24] K. Maity, B. Mahanty, T.K. Sinha, S. Garain, A. Biswas, S.K. Ghosh, S. Manna, S.K. Ray, D. Mandal, Two-Dimensional piezoelectric MoS2-modulated nanogenerator and nanosensor made of poly(vinlydine fluoride) nanofiber webs for selfpowered electronics and robotics, Energy Technol. 5 (2017) 234e243. [25] A. Sultana, P. Sadhukhan, M.M. Alam, S. Das, T.R. Middya, D. Mandal, OrganoLead halide perovskite induced electroactive b-Phase in porous PVDF films: an excellent material for photoactive piezoelectric energy harvester and photodetector, ACS Appl. Mater. Interfaces 10 (2018) 4121e4130. [26] A. Tamang, S.K. Ghosh, S. Garain, M.M. Alam, A.J. Haeberle, K. Henkel, D. Schmeisser, D. Mandal, DNA-assisted b-phase nucleation and alignment of molecular dipoles in PVDF film: a realization of self-poled bioinspired flexible polymer nanogenerator for portable electronic devices, ACS Appl. Mater. Interfaces 7 (2015) 16143e16147. [27] R. Ozerov, A. Vorobyev, in: Physics for Chemists, Elsevier Science, 2007, p. 251. [28] T. Ikeda, in: Fundamentals of Piezoelectricity, Oxford University Press, Oxford, 1996, p. 1. [29] A. Erturk, D.J. Inman, Piezoelectric Energy Harvesting, John Wiley & Sons, New York, 2011. [30] F. Wang, M. Tanaka, S. Chonan, Development of a PVDF piezopolymer sensor for unconstrained in-sleep cardiorespiratory monitoring, J. Intell. Mater. Syst. Struct. 14 (2003) 185e190. [31] A. Manbachi, R. Cobbold, Development and application of piezoelectric materials for ultrasound generation and detection, Ultrasound 19 (2011) 187e196. [32] W.P. Mason, Piezoelectricity, its history and applications, J. Acoust. Soc. Am. 70 (1981) 1561e1566. [33] G. Zhu, A. Wang, Y. Liu, Y. Zhou, Z. Wang, Functional electrical stimulation by nanogenerator with 58 V output voltage, Nano Lett. 12 (2012) 3086e3090. [34] H. Kim, J. Kim, J. Kim, A review of piezoelectric energy harvesting based on vibration, Int. J. Precis. Eng. Manuf. 12 (2011) 1129e1141. [35] T. Yucel, P. Cebe, D. Kaplan, Structural origins of silk piezoelectricity, Adv. Funct. Mater. 21 (2011) 779. [36] C. Wan, C.R. Bowen, Multiscale-structuring of polyvinylidene fluoride for energy harvesting: the impact of molecular-, micro- and macro-structure, J. Mater. Chem. A 5 (2017) 3091e3128. [37] H.S. Nalwa, Ferroelectric Polymers: Chemistry: Physics, and Applications, CRC Press, 1995. [38] P. Martins, A.C. Lopes, S. Lanceros-Mendez, Electroactive phases of poly(vinylidene fluoride): determination, processing and applications, Prog. Polym. Sci. 39 (2014) 683e706.

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[73] D. Mandal, K.J. Kim, J.S. Lee, Simple synthesis of palladium nanoparticles, b-phase formation, and the control of chain and dipole orientations in palladium-doped poly(vinylidene-fluoride) thin films, Langmuir 28 (2012) 10217e10310. [74] D. Mandal, K. Henkel, D. Schmeißer, The electroactive b-phase formation in Poly(vinylidene fluoride) by gold nanoparticles doping, Mater. Lett. 73 (2012) 123e125. [75] X. Liu, J. Ma, X. Wu, L. Lin, X. Wang, Polymeric nanofibers with ultrahigh piezoelectricity via self-orientation of nanocrystals, ACS Nano 11 (2017) 1901e1910. [76] A. Batth, A. Mueller, L. Rakesh, A. Mellinger, in: Presented at the 2012 Annual Report Conference on Electrical Insulation and Dielectric phenomena(CEIDP), Montreal, QC, Canada, 2012, pp. 28e31. [77] L. Wu, W. Yuan, N. Hu, Z. Wang, C. Chen, J. Qiu, J. Ying, Y. Li, Improved piezoelectricity of PVDF-HFP/carbon black composite films, J. Phys. D Appl. Phys. 47 (2014) 135302. [78] P. Adhikary, S. Garain, D. Mandal, The co-operative performance of a hydrated salt assisted sponge like P(VDF-HFP) piezoelectric generator: an effective piezoelectric based energy harvester, Phys. Chem. Chem. Phys. 17 (2015) 7275e7281. [79] P. Adhikary, S. Garain, S. Ram, D. Mandal, Flexible hybrid Eu3þ doped P(VDFHFP) nanocomposite film possess hypersensitive electronic transitions and piezoelectric throughput, J. Polym. Sci., Part B Polym. Phys. 54 (2016) 2335. [80] J.S. Dodds, F.N. Meyers, K.J. Loh, Piezoelectric characterization of PVDF-TrFE thin films enhanced with ZnO nanoparticles, IEEE Sens. J. 1889 (2012) 12. [81] C.-G. Wu, G.-Q. Cai, W.-B. Luo, Q.-X. Peng, X.-Y. Sun, W.-L. Zhang, Enhanced pyroelectric properties of PZT/PVDF-TrFE composites using calcined PZT ceramic powders, J. Adv. Dielectr. 03 (2013) 1350004. [82] A.K. Batra, M.E. Edwards, A. Alomari, A. Elkhaldy, Am. J. Mater. Sci. 5 (2015) 55. [83] S.K. Ghosh, K. Roy, H.K. Mishra, M.R. Sahoo, B. Mahanty, P.N. Vishwakarma, D. Mandal, Rollable magnetoelectric energy harvester as wireless IoT sensor, ACS Sustain. Chem. Eng. 8 (2020) 864e873. [84] W.K. Sakamoto, S. Shibatta-Kagesawa, D.H.F. Kanda, S.H. Fernandes, E. Longo, G.O. Chierice, Piezoelectric effect in composite polyurethaneeferroelectric ceramics, Phys. Status Solidi A. 172 (1999) 265. [85] C. David, J.F. Capsal, L. Laffont, E. Dantras, C. Lacabanne, Piezoelectric properties of polyamide 11/NaNbO3 nanowire composites, J. Phys. D Appl. Phys. 45 (2012) 415305. [86] H. Khanbareh, S. Van Der Zwaag, W.A. Groen, Piezoelectric and pyroelectric properties of conductive polyethylene oxide-lead titanate composites, Smart Mater. Struct. 24 (2015) 045020. [87] A. Sultana, M.M. Alam, A. Biswas, T.R. Middya, D. Mandal, Fabrication of wearable semiconducting piezoelectric nanogenerator made with electrospun-derived zinc sulfide nanorods and poly(vinyl alcohol) nanofibers, Transl. Mater. Res. 3 (2016) 045001. [88] G. Zhao, X. Zhang, X. Cui, S. Wang, Z. Liu, L. Deng, A. Qi, X. Qiao, L. Li, C. Pan, Y. Zhang, L. Li, Piezoelectric polyacrylonitrile nanofiber film-based dual-function self-powered flexible sensor, ACS Appl. Mater. Interfaces 10 (2018) 15855. [89] D. Dhakras, S. Ogale, High-performance organiceinorganic hybrid piezonanogenerator via interface enhanced polarization effects for self-powered electronic systems, Adv. Mater. Interfaces (2016) 1600492. [90] P. Adhikary, D. Mandal, Enhanced electro-active phase in a luminescent P(VDFe HFP)/Zn2þ flexible composite film for piezoelectric based energy harvesting applications and self-powered UV light detection, Phys. Chem. Chem. Phys. 19 (2017) 17789e17798.

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[91] R. Pérez, M. Kral, H. Bleuler, Study of polyvinylidene fluoride (PVDF) based bimorph actuators for laser scanning actuation at kHz frequency range, Sensor. Actuator. 183 (2012) 84e94. [92] X. Gong, M. Tong, Y. Xia, W. Cai, J.S. Moon, Y. Cao, G. Yu, C.L. Shieh, B. Nilsson, A.J. Heeger, High-detectivity polymer photodetectors with spectral response from 300 nm to 1450 nm, Science 325 (2009) 1665e1667. [93] T.K. Sinha, S.K. Ghosh, R. Maiti, S. Jana, B. Adhikari, D. Mandal, S.K. Ray, Graphene-silver-induced self-polarized PVDF-based flexible plasmonic nanogenerator toward the realization for new class of self-powered optical sensor, ACS Appl. Mater. Interfaces 8 (2016) 14986e14993. [94] R. Costa, C. Ribeiro, A.C. Lopes, P. Martins, V. Sencadas, R. Soares, S. LancerosMéndez, Osteoblast, fibroblast and in vivo bio-logical response to poly(vinylidene fluoride) based composite materials, J. Mater. Sci. Mater. Med. 24 (2012) 395e403. [95] C. Ribeiro, S. Moreira, V. Correia, V. Sencadas, J.G. Rocha, F.M. Gama, J.L.G. Ribelles, S. Lanceros-Méndez, Enhanced proliferation of pre-osteoblastic cells by dynamic piezoelectric stimulation, RSC Adv. 2 (2012) 11504e11509. [96] V.K. Sharma, R.A. Yngard, Y. Lin, Silver nanoparticles: green synthesis and their antimicrobial activities, Adv. Colloid Interface Sci. 145 (2009) 83e96. [97] A.C. Lopes, C. Caparros, J.L. Ribelles, I.C. Neves, S. Lanceros-Méndez, Electrical and thermal behavior of gamma-phase poly(vinylidenefluoride)/NaY zeolite composites, Microporous Mesoporous Mater. 161 (2012) 98e105. [98] P. Martins, X. Moya, L.C. Phillips, S. Kar-Narayan, N.D. Mathur, S. LancerosMéndez, Linear anhysteretic direct magnetoelectric effect in Ni0.5Zn0.5Fe2O4/poly(vinylidene fluoride-trifluoroethylene) 0e3nanocomposites, J. Phys. D 44 (2011), 482001/1-4. [99] P. Martins, A. Lasheras, J. Gutierrez, J.M. Barandiaran, I. Orue, S. Lanceros-Méndez, Optimizing piezoelectric and magnetoelectric responseson CoFe2O4/P(VDF-TrFE) nanocomposites, J. Phys. D 44 (2011), 495303/1-7. [100] S.M. Damaraju, S. Wu, M. Jaffe, T.L. Arinzeh, Structural changes in PVDF fibers due to electrospinning and its effect on biological function, Biomed. Mater. 8 (2013) 045007. [101] T. Marques-Almeida, V.F. Cardoso, S. Ribeiro, F.M. Gamad, C. Ribeiro, S. Lanceros-Mendez, Tuning myoblast and pre-osteoblast cell adhesion site, orientation and elongation through electroactive micropatterned scaffolds, ACS Appl. Bio Mater. 2 (2019) 1591e1602.

CHAPTER EIGHT

Advances of electrochromic and electro-rheological materials R. Alves1, 2, M.M. Silva3 1

Center of Physics, University of Minho, Campus de Gualtar, Braga, Portugal Institute of Science and Innovation for Bio-Sustaninability (IB-S), University of Minho, Campus de Gualtar, Braga, Portugal 3 Department/Center of Chemistry, University of Minho, Campus de Gualtar, Braga, Portugal 2

8.1 Advances of electrochromic materials Chromogenics are a family of materials where the transmittance of visible light and solar energy is to be varied under an external action [1]. The optical change of the materials can be induced by different stimuli leading to thermochromic, photochromic, and electrochromic (EC) materials (Fig. 8.1). “Thermochromism” and “photochromism” are effects described as changes of color produced by heat or light, while the shift in the absorption and emission spectra upon application of an electric field is an effect called as “electrochromism”. Thermochromism can be defined as a thermally induced reversible color change. There are different thermochromic materials, such as organic and inorganic compounds [2]. Thermochromic liquid crystals are important materials that show different colors at different temperatures due to the selective reflection of specific wavelengths of light from their structures [3]. Organic thermochromic mixtures are another category of materials used to impart colors that change with temperature that involves organic dyes [3]. The thermochromic behavior can also be found in spiropyrans, thermochromic ethylenes, sulfur compounds, and miscellaneous compounds [4]. Thermochromic window coatings offer some possibilities to achieve energy efficiency in windows, and according to this approach, a single film is all that is needed, and the most widely studied thermochromic material is VO2 (vanadium dioxide); however, the benefits are not as large as for electrochromics [5]. A thermochromic solar cell utilizing the structural phase transitions in inorganic halide perovskite cesium lead iodide/bromide for smart photovoltaic window applications demonstrate high thermal stability and fully reversible color and performance [6]. Thermochromic Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00003-2

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Chromogenics

Thermocromics

Photochromics

Electrochromics

Applications:

Applications:

Food packaging

Sunglasses

Smart windows

Babies cutlery

Toys

Rear view mirrors

Cups

Clothing

Protector eyewear

Contact thermometers

Cosmetics

Sunglasses

Applications:

Figure 8.1 The different types of chromogenic materials.

material 2,4,5-triphenylimidazole (also called lophine) can be also used for the fabrication of temperature optical fiber sensors for monitoring the temperature of the water in several applications [2]. Photochromism can be simply defined as the light-induced reversible color change. It is a term used for a reversible transformation of a specific molecule induced by different wavelengths of light, where several physical and chemical properties may be tuned by light [7]. In most photochromic compounds the stable form is colorless or pale yellow, and when irradiated they acquire coloration (positive photochromism), while less common photochromic compounds present a colored form, and when irradiated a colorless form is observed (negative or inverse photochromism), or exhibit a reversible change between different colors [8]. Photochromic materials are very promising and have potential in several scientific research fields, such as chemistry, physics, materials science, biology, and nanotechnology [9]. Photochromic materials have been widely investigated and can be divided into organic, inorganic, and organic-inorganic hybrid materials. Some families of organic compounds that show photochromism, such as, spiropyrans, spirooxazines, chromenes, fulgides and fulgimides, diarylethenes and related compounds, spirodihydroindolizines, azo compounds, polycyclic aromatic compounds, anils and related compounds (hydrogen transfer), polycyclic quinones (periaryloxyquinones), perimidinespirocyclohexadienones, viologens, and triarylmethanes [10]. In comparison with photochromic organic materials, inorganic photochromic materials, such as transition metal oxides, are extensively studied due to their good stability and cost efficiency, and tungsten oxide (WO3) has attracted considerable attention [11]. Generally, there are a variety of organic compounds, and

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they are easy to modify or process, while inorganic compounds have considerable thermal stability, high strength, and diverse coordination chemistry. In organic-inorganic hybrid photochromic materials, the distinct synergetic characteristics between each component offer various possibilities to obtain new “smart”, high performing, and tailor-made photochromic materials because in hybrid materials it is possible to preserve or even improve the respective features of the components and also produce new properties depending on the synergy between them [12]. The electrochromism is defined as a physicochemical phenomenon where reversible and visible changes in transmittance and/or reflectance are associated to electrochemically induced oxidation-reduction reaction, and the color change is between a transparent state (called “bleached” state) and a colored state, or between two colored states [13]. There are three main types of EC materials, Fig. 8.2, which may be divided according to their possible optical states [14]. The first includes metal oxides, viologens, and polymers such as poly(3,4-ethylenedioxythiophene) (PEDOT), which are materials with one colored and one bleached state. They are very useful for absorption/transmission-type devices such as “smart windows” and optical shutters. The second type is related to materials with two distinctive colored states, which are useful for display applications where different colors in different redox states are desired. A good example of this type is polythiophene that switches from red to blue upon oxidation. The third and last one includes electrochromes where more than two color states are possible depending on the redox state of the material. Conjugated polymers are inherently multicolor materials that are interesting due to their versatility for making copolymers, laminates, and blends. Polyaniline (PANI)

EC materials

One colored state

Two colored states

More than two color states

Examples: Metal oxides

Examples:

Examples:

Viologens

Polythiophene

Conjugated polymers

PEDOT

Figure 8.2 Classification of the EC materials according to the possible colored states.

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or poly(3,4-propylenedioxypyrrole) (PProDOP) are examples of such polymers. All these types of EC materials will be discussed in detail in the sections below. Electrochromic contrast, coloration efficiency, switching speed, stability, and optical memory are important parameter needs to evaluate the performance of an EC material, Fig. 8.3. The electrochromic contrast, or optical contrast, is described in terms of the percent transmittance change (DT (%) ¼ Tbleached  Tcolored) at a selected wavelength where the material has high optical contrast [15]. The optical contrast quantifies how well the electrochromic compound is turned “on” and “off” and so it is a key factor for evaluating an EC material [16]. The coloration efficiency (CE), also denoted as the electrochromic efficiency, determines the amount of charge necessary to produce the optical change, usually represented in cm2/C (area per charge) according to Eq. (8.1), CE ¼

DðODÞ DQ

(8.1)

where D(OD) is the optical density change (D(OD) ¼ log (Tcolored/ Tbleached)) and DQ is the electronic charge [15]. Switching speed, also referred as response time, is the time required for the coloring/bleaching process, and it is obtained by correlating the change

EC performance

Electrochromic contrast

Coloration efficiency

Switching speed

Electrochromic stability

Optical memory

Figure 8.3 Some important parameter needs to evaluate the performance of EC materials.

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in transmittance with time required for the change, in response to an applied potential. The response time is dependent of many factors such as magnitude of the applied potential, morphology of thin film, film thickness, ion diffusion in thin films, and conductivity of the electrolyte. Therefore, switching speed is based on the changes in optical density, and faster switching speed is important [17]. Electrochromic stability is associated with electrochemical stability and a main drawback is related to the stabilities upon cycles, and lack of durability. Cycle stability testing provides important information regarding the viability of device components in commercial applications [18]. The optical memory, also called open-circuit memory or memory effect, is the time that the material retains its absorption state after the electric field is removed; in other words, it is the ability to maintain a given redox form when taken to open-circuit, allowing to maintain a color state without a continuous power supply [16,19]. The memory effect is dependent on the chemical environment surrounding the electrochromic polymer film, namely, the counter species and the diffusion of the charge carriers (affected by electrolyte composition, viscosity, and thickness) [16]. In the last decades, the growing interest in EC materials is due to their potential applications, such as, in “smart” windows [20,21], high-quality paperlike displays [22,23], visors for motorcycle helmets [24,25], and smart sunglasses [26], for example.

8.1.1 Materials for electrochromic devices construction EC materials have been extensively studied due to their potential applications, as in ECDs, and electrochemical techniques, such as cyclic voltammetry, coulometry, and chronoamperometry, are appropriate measurements for their characterization. The ECDs are considered a promising technology for next-generation electronic applications, and they operate essentially as a rechargeable battery, where the color change is associated with the charging and discharging of the cell with applied potential of a few volts [27]. There are many configurations of ECDs with layers of different compositions, prepared by various techniques, and Fig. 8.4 shows a typically sandwich multilayer configuration prototype composed by a glass substrate, a transparent conductor, an ion storage coating, an ionic conductor, and an electrochromic coating. The ECD presents a multicomponent structure, and one of them can influence the device performance [28].

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ro

c

o hr

m

ic

an Tr

m fil

sp

a

n re

tc

o

u nd

G

la

ct

or

ss

– – –

+ + +

– – –

+ + +

– – –

Ions

+

–

Figure 8.4 Schematic illustration of an ECD showing the movement of ions/electrons under an externally applied electric field. Reprinted by permission from C.G. Granqvist, Out of a niche, Nat. Mater. 5 (2) (2006) 89e90, https://doi.org/10.1038/nmat1577, Copyright (2006) Nature.

8.1.1.1 Optical transparent electrodes The transparent conductors (TCs) appear as the top and bottom electrodes. Fluorine tin oxide (FTO) is one common choice, which comprises FTO on glass. It presents desirable properties such as good transparent conducting performance, low cost, and good chemical durability, which make it an interesting material used as a transparent conducting layer for the ECDs or dye-sensitized solar cell [29]. FTO electrodes can be fabricated using the horizontal ultrasonic spray pyrolysis deposition, for example [30]. Indium tin oxide (ITO) is one of the most commonly used material that appears sputtered on glass as a thin film. ITO shows high transmittance (>80%) at visible wavelengths, and low sheet resistance ( Y 2

Y1

]

X ---> Y* Y2

(b) Combination effect

Phase 1

Phase 1 : X ---> Y 1/Z1 Phase 2 : X ---> Y 2/Z2

]

Phase 2

X ---> (Y/Z)* Improvement

Y1 Y2

Phase 1

Phase 2

Y1/Z1

Y2/Z2

Z1

Phase 1

Z2 Phase 1

Phase 2

Phase 2

(c) Product effect Phase 1 : X ---> Y Phase 2 : Y ---> Z

]

X ---> Z

New function

Figure 10.6 Composite effects: sum, combination, and product effects [35].

A dramatic enhancement in the mechanical strength of the rod is achieved by adding carbon fibers in a special orientation, i.e., along a rod (showing a convex relation as depicted in Fig. 10.6(a)). Another interesting example is an NTC-PTC material [36]. V2O3 powders are mixed in epoxy with a relatively high packing rate (3-3), as illustrated in Fig. 10.7. Since V2O3 exhibits a semiconductor-metal phase transition at 113 C, a drastic resistivity change is observed with increasing temperature. A further increase in temperature results in a larger thermal expansion for epoxy than for the ceramic, leading to a separation of each particle, and the structure becomes a 0-3 composite. The V2O3 particle separation increases the resistivity significantly at around 100 C. Thus, the conductivity of this composite is rather high only over a limited temperature range (around 110 to 100 C), which is sometimes called the “conductivity window.” 10.2.2.2 Combination effects In certain cases, the averaged value of the output, Y*, of a composite does exceed Y1 and Y2. This enhanced output refers to an effect on a figure-of-

Kenji Uchino

Resistivity (Ω.cm)

368

106 104 102 100 –100

0

100

200

Temperature (°C)

Figure 10.7 NTC-PTC effect observed in a V2O3: epoxy composite [36].

merit Y/Z which depends on two parameters Y and Z. Suppose the parameters Y and Z follow concave and convex type sum effects, respectively, as illustrated in Fig. 10.6(b), the combination value Y/Z will exhibit a maximum at an intermediate ratio of phases. This is called a “combination effect.” Certain piezoelectric ceramic:polymer composites exhibit a combination property of g (the piezoelectric voltage constant), which is provided by d/ε (d: piezoelectric strain constant, and ε: permittivity). In the 1-3 piezoelectric composite, where PZT rods are arranged in silicone rubber matrix [37], the effective piezoelectric constant d*33 of the composite is rather close to that of the PZT rod itself d33 down to 10 volume % of PZT, while the effective permittivity is almost proportional to the volume % of PZT. Thus, we can realize 10 times higher piezoelectric constant g*33 in a 10 vol % PZT composite, in comparison with the pure PZT specimen. 10.2.2.3 Product effects When Phase 1 exhibits an output Y with an input X, and Phase 2 exhibits an output Z with an input Y, we can expect for the composite an output Z with an input X. A completely new function is created for the composite structure, called a “product effect.” [See Fig. 10.6(c).] Functionality matrix

The author introduces his “Functionality Matrix” concept here. Table 10.3 lists the various effects in various materials/devices, relating the input (electric field, magnetic field, stress, heat and light) with the output (charge/ current, magnetization, strain, temperature, and light). Electrically conducting and elastic materials, which generate current and strain outputs,

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respectively, under the input, voltage, or stress, are sometimes called “trivial” materials (well-known diagonal coupling phenomena, that is, Ohm’s and Hooke’s Laws). On the other hand, pyroelectric and piezoelectric materials, which generate an electric field with the input of heat and stress (unexpected phenomena!), respectively, are called “smart” materials. These off-diagonal couplings have corresponding “converse effects,” the electrocaloric and converse-piezoelectric effects, and both “sensing” and “actuating” functions can be realized in the same materials. Pure ferroelectric materials exhibit most of these effects with the exception of the magnetic phenomena. Thus, ferroelectrics are said to be very “smart” materials. Taking into account ð5 5Þ components of Table 10.3, we introduce a ð5 5Þ “functionality matrix.” We invented “photostrictive” actuators, according to the following development procedure [38]. If one material has “photovoltatic effect,” the functionality matrix of this material can be expressed by: 1 0 0 0 0 0 Light Emission C B 0 0 0 0 0 C B C B C: B 0 0 0 0 0 C B C B 0 0 0 0 0 A @ Photovoltaic 0 0 0 0

Table 10.3 Various effects in materials/devices.

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On the other hand, a piezoelectric has a functionality matrix of the following form: 1 0 0 0 Converse Piezo. 0 0 C B 0 0 0 0 0C B C B B Piezo Ele. 0 0 0 0C C: B C B 0 0 0 0 0A @ 0

0

0

0

0

As a composite effect of photovoltaic and piezoelectric, when the light illumination is input first, the expected phenomenon is expressed by the matrix product: 1 0 0 0 0 0 Light Emis. C B C B 0 0 0 0 0 C B C B C5 B 0 0 0 0 0 C B C B C B 0 0 0 0 0 A @ Photovoltaic 0 0 0 0 1 0 0 0 Converse Piezo 0 0 C B 0 0 0 0 0C B C B C B Piezo Ele. 0 0 0 0 C B C B 0 0 0 0 0A @ 0 0 0 0 0 1 0 0 0 0 0 0 C B 0 0 0C B0 0 C B 0 0 0C ¼B C: B0 0 C B 0 0 0A @0 0 0 0 Photostriction 0 0 Note that only one component, “photostriction” is derived from this product calculation. Now, let us consider “magnetoelectric” effect. If one material has “piezomagnetic effect” and its converse effect “magnetostrictive effect,” the functionality matrix of this material is expressed as follows:

Magnetoelectric composite materials: a research and development case study

0

0

B B0 B B0 B B @0 0

0 0

0

0

Mag. Strictive 0

Piezo. Mag

0

0

0

0

0

0

0

0

0

371

1

C 0C C 0C C: C 0A 0

On the other hand, a piezoelectric has a functionality matrix: 1 0 0 0 Converse Piezo. 0 0 C B 0 0 0 0 0C B C B B Piezo Ele. 0 0 0 0C C: B C B 0 0 0 0 0A @ 0 0 0 0 0 Thus, in a composite composed of the above materials, when the magnetic field is input first, the expected phenomenon is expressed by the matrix product: 1 0 0 0 0 0 0 C B C B 0 Mag. str. 0 0 C B0 C B C B C5 B 0 Piezo. Mag 0 0 0 C B C B C B 0 0 0 0C B0 A @ 0 0 0 0 0 1 0 0 0 Converse Piezo 0 0 C B C B 0 0 0 0 0C B C B C B B Piezo Ele. 0 0 0 0C C B C B C B 0 0 0 0 0C B A @ 0 0 0 0 0 0 1 0 0 0 0 0 B C B C Mag. Ele. 0 0 0 0 B C B C B C B ¼B 0 0 0 0 0C C: C B C B 0 0 0 0 0C B A @ 0 0 0 0 0

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Note that only one component, “magnetoelectric coupling” is derived from this product calculation. To the contrary, if we start from the electric field input first, the expected phenomenon will be: 1 0 0 0 Converse Piezo 0 0 C B C B 0 0 0 0 0C B C B C B B Piezo Ele. 0 0 0 0C C5 B C B C B 0 0 0 0 0C B A @ 0 0 0 0 0 1 0 0 0 0 0 0 C B C B 0 Mag. str. 0 0 C B0 C B C B C B 0 Piezo. Mag 0 0 0 C B C B C B 0 0 0 0 0 C B A @ 0 0 0 0 0 1 0 0 Ele. Mag. 0 0 0 C B C B 0 0 0 0C B0 C B C B B 0 0 0 0C ¼ B0 C: C B C B 0 0 0 0C B0 A @ 0 0 0 0 0 Note now that the resulting product matrix includes only one component, electromagnetic effect. Magnetoelectric 0-0 composites

Philips Laboratories in Netherlands developed ME materials based on the above concept [35,39]. The material was composed of magnetostrictive CoFe2O4 and piezoelectric BaTiO3 mixed and sintered together. Fig. 10.8(a) shows a micrograph of a transverse section of a unidirectionally solidified rod of the materials with an excess of TiO2 (1.5 wt.%). The fourfinned spinel dendrites CoFe2O4 and white cubical barium titanate grains

Magnetoelectric composite materials: a research and development case study

Hac (100e at 1 kHz)

1 mm

(mV/cm.Oe)

(b)

dE — dH

(a)

373

150 100 50

0

0.5

1.0 1.5

2.0

Hdc (kOe)

Figure 10.8 (a) Micrograph of a transverse section of a uni-directionally solidified rod of mixture of magnetostrictive CoFe2O4 and piezoelectric BaTiO3, with an excess of TiO2. (b) Magnetic field dependence of the magnetoelectric effect in a CoFe2O4: BaTiO3 composite (at room temperature) [35,39]. Note maximum (dE/dH) z 130 mV/cm∙Oe.

are observed in cells (100). Ideally CoFe2O4 and BaTiO3 grains should be 0-0 connectivity, without connecting each phase in a long range. Fig. 10.8(b) shows the magnetic field dependence of the magnetoelectric effect in a CoFe2O4: BaTiO3 composite (at room temperature). Note maximum (dE/dH) z 130 mV/cm∙Oe, which is more than 10 times higher than the typical values observed in single-phase ME materials. In parallel to other research groups, the author’s team also chased initially the 0-0 composites by using PZT powders (rather than barium titanate) for enhancing the magnetoelectric performance [40]. We investigated intensively the effect of the sintering temperature on the sintering behavior, microstructures, piezoelectric and ME properties of this ME particulate composite with Ni-ferrite doped with Co, Cu, Mn particles, and PZT matrix [40]. Not only the connectivity of the ferrite phase, but also the sintering temperature are the important parameters for higher ME voltage coefficient (dE/dH). The chemical reaction of the PZT with ferrite, and connection of ferrite particles make it difficult to get high ME effects. We obtained the highest ME voltage coefficient from the composite with 20% ferrite added and sintered at 1250 C. A homogeneous and well-dispersed microstructure, no chemical reaction between the two phases, and large grain size of the matrix PZT phase were the most important factors to get a high ME voltage coefficient. Though our value was 45% higher than the previously reported value from Philips Lab., the

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enhancement was not so significant as our large efforts. Thus, we basically abandoned the “particulate” composite approach, and the target was shifted to the “laminated” composites.

10.3 Designing of ME laminate composites We describe how to optimize the design of ME laminate composites in this section.

10.3.1 Challenge to the laminate composites The particulate ME composites (0-0 type) made of piezoelectric and magnetostrictive ferrite materials showed higher ME properties compared with single-phase ME materials such as Cr2O3, as we mentioned in the previous section. However, these composites still need some important issues addressed when fabricating the sintered ME particulate composites to obtain superior ME response. First, no chemical reaction should occur between the piezoelectric and magnetostrictive materials during the sintering process. The chemical reaction may reduce the piezoelectric or magnetostrictive properties of each phase. Second, the resistivity of the magnetostrictive phase should be as high as possible. If the resistivity of the magnetostrictive phase is low, the electric poling of the piezo-phase becomes very difficult due to leakage current. Also, the leakage reduces the ME output voltage of the composites. When the ferrite particles make connected chains (2D, 3D), the electric resistivity of the composites is reduced significantly because of the low resistivity of the ferrite. Therefore, good dispersion of the ferrite particles in the piezoelectric matrix is required in order to sustain sufficient electric resistivity of the composite. Third, mechanical defects such as pores at the interface between the two phases should not exist in the composite for good mechanical coupling, through which magnetostrictive and piezoelectric materials interact. These difficulties may be overcome by using a laminar composite (2-2 type), because no chemical reactions and dispersions are involved in the fabrication process. In addition to these advantages, the laminated ME composites have a very simple structure and relatively simple fabrication method, i.e., bonding each disk. The laminated ME composites can be easily applied to practical applications, such as magnetic field sensing devices, leak detectors for microwave ovens, and current measurement of high-power

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electric transmission systems. Thus, we invented the laminated ME composites made by using piezoelectric and magnetostrictive materials [41e44]. Lead Zirconate Titanate (PZT) and Terfenol-D disks were used as the piezoelectric material and magnetostrictive material, respectively, in the ME laminate composites. The composites were manufactured by sandwiching and bonding a PZT disk between two layers of Terfenol-D disks, as shown in Fig. 10.9. When a magnetic field is applied to this composite, the top and bottom Terfenol-D disks shrink or expand. This shrinkage or expansion generates stresses in the sandwiched piezoelectric PZT disk. Hence, electric signals can be obtained when the composite is subjected to a magnetic field. To optimize the PZT and Terfenol-D ME laminate composites, we investigated [1] the effect of the piezoelectric properties of the piezoelectric layer (PZT and PMN-PT single crystals) [2], thickness ratio between the PZT and Terfenol-D disks [3], the directional dependence of the magnetostriction of the Terfenol-D disks and [4] of the a.c. magnetic field on the ME response of the PZT/Terfenol-D laminate composites.

(a) Polarization Terfenol-D PZT

Magnetic field

(b) Permanent Magnet

Terfenol-D PZT 12.7 mm

Figure 10.9 ME laminate composite using Terfenol-D and PZT disks. (a) schematic structure and (b) photograph of the device.

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10.3.2 ME measuring technique Prior to the device designing principle, we introduce the ME measuring technique [41]. The ME voltage coefficient was determined by measuring the electric voltage (electric field) generated across the sample when an a.c. magnetic field with a d.c. bias was applied to it. The magnetoelectric property was measured in terms of the variation of the coefficient dE/dH as a function of d.c. magnetic bias field. An electromagnet (GMW 5403 Magnet, Power and Buckley Inc., New Zealand) was used for the bias magnetic flux density B up to 0.45 T (i.e., magnetic field H ¼ 4.5 kOe). The frequency dependence of the magnetoelectric voltage coefficient of composites was determined in the range of 10 Hz to 3 kHz under a 1 kOe d.c. bias. The coefficient was measured directly as response of the sample to an a.c. magnetic input signal at 1 kHz and 2 Oe amplitude (Helmholtz coils were used to give a uniform a.c. magnetic field in the space between the coils) superimposed on the d.c. bias field, both parallel to the sample axis (Fig. 10.10). A signal generator (33120A, Hewlett Packard Co., USA) was used to drive the Helmotz coils and generate the a.c. magnetic field. The voltage generated in the piezoelectric layer was measured under an open circuit condition. A differential amplifier based on the INA121 FET-input Instrumentation Amplifier (Burn-Brown Inc.) was used. This amplifier is specially designed for high impedance transducers, providing differential input impedance in the order of 1012GU/1 pF, which represents almost an ideal open circuit condition. The electric circuit of the amplifier is shown in Fig. 10.10(c). The output signal from the amplifier was measured with an oscilloscope (54645A, Hewlett Packard Co., USA). The output voltage divided by the thickness of the sample and the a.c. magnetic field gives the magnetoelectric voltage coefficient of the samples.

10.3.3 Effect of piezoelectric properties 10.3.3.1 PZT case The effect of the piezoelectric properties of the PZT material on the ME response of the laminate composites was studied first on samples consisting of two Terfenol-D disks and one different kind of PZT materials. Note that the permanent magnets in Fig. 10.9(b) were not attached for the basic measurement, but they were used for a practical application to provide the optimized d.c. bias magnetic field to the ME device. Table 10.4 shows the piezoelectric properties of the PZT materials used for this study. APC 840,

Magnetoelectric composite materials: a research and development case study

(a)

377

Electromagnet

Helmholtz coils

Hall probe Sample Aluminum shield

Charge amplifier

Function generator

DC power supply

Digitial teslameter Oscilloscope

(b)

V+

(c)

0.1 PF

Over voltage protection

40k:

40k:

– VIN

25k:

–

25k:



R

– Over voltage protection

VOUT

40k:



40k:

INA121 0:

0.1 PF

0:

V–

Figure 10.10 (a) Schematic diagram of ME measurement system, (b) Photo of the measurement system, (c) Amplifier circuit for magnetoelectric voltage measurement.

Kenji Uchino

378 Table 10.4 Piezoelectric properties of PZT materials used in this study. Material εT33 /ε0 tand (%) d33(pC/N) g33(mVm/N) Qm

kp

APC840 PZT-5A APC841

0.59 0.57 0.60

1250 1730 1250

0.4 1.5 0.35

320 340 275

500 68 1400

25.6 19.6 22.0

Bold values repesents the Top data among three PZT’s, which indicates the best performance among three materials in terms of that particular figure-of-merit.

PZT-5A, and APC 841 were used for their high g33, high d33, and high Qm, respectively [41]. Fig. 10.11 shows the ME voltage coefficient variation as a function of the d.c. magnetic bias with three different PZT types. All PZT disks were machined down to the same thickness (0.5 mm) and diameter (12.7 mm). The ME voltage coefficients of all the composites were increased with increasing d.c. bias until saturated around 4 kOe. Since the sensitivity is mainly determined by the piezoelectric voltage coefficient (g33), the composite with APC-840 PZT showed the most superior ME property. The maximum ME voltage coefficient for this composite was 4.68 V/cm$Oe under 4.2 kOe d.c. magnetic bias or higher. Note that this 2-2 type laminate composite exhibits more than 30 times higher ME voltage coefficient than the 0-0 type particulate composite introduced in Fig. 10.8. The dependence of the magnetoelectric voltage coefficient of composites on the frequency of a.c. magnetic field was observed under a fixed d.c. magnetic bias at 1 kOe for all samples, and found a maximum

dE/dH (mV/Oe cm)

5000 APC 840 PZT-5A APC 841

4000 3000 2000 1000 0

0

1000

2000

3000

4000

5000

Magnetic field (Oe)

Figure 10.11 ME voltage coefficient as a function of applied d.c. magnetic bias for various PZT disks (APC 840 for high g33, PZT-5A for high d33, APC 841 for high Qm) at 1 kHz [41].

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magnetoelectric voltage coefficient at around 150 Hz. This frequency dependence seems to occur due to a mechanical resonance of the sample device and sample holder, and is not the property of the composite itself. Since the lowest mechanical resonance frequency of these composite samples can be evaluated to be more than 100 kHz, a significant frequency dependence cannot be expected in this frequency range. 10.3.3.2 PMN-PT single crystal case As shown in the previous section, the most important factors needed in order to achieve a high ME voltage coefficient from piezoelectricmagnetostrictive composites are a high piezoelectric voltage coefficient (gij). Since our discovery relaxor single crystals such as Pb(Mg1/3Nb2/3) O3-PbTiO3 (PMN-PT) and Pb(Zn1/3Nb2/3)O3-PbTiO3 (PZN-PT) are very well known to have superior piezoelectric properties [45,46], extremely high ME coefficients are expected, when the PZT ceramic layer is replaced by a (001) oriented Pb(Mg1/3Nb2/3)O3-PbTiO3 (PMN-PT) single crystal, which has a much higher gij coefficient. Fig. 10.12 shows the ME voltage coefficient as a function of applied d.c. magnetic field. The magnetic field dependence of dE/dH was similar for all three types of piezoelectric specimens. As is evident in the figure, the ME voltage coefficient increased with increasing magnetic bias, saturating at a bias level of w4 kOe. The ME laminate composite made using a PMN-PT single crystal had the highest ME voltage coefficient. The value of dE/dH was

12 PMN-PT single crystal PZT ceramic PMN-PT ceramic

dE/dH (V/cm Oe)

10 8 6 4 2 0 0

1000

2000

3000

4000

5000

Magnetic bias (Oe)

Figure 10.12 ME voltage coefficient for laminate composites as a function of an applied d.c. magnetic bias field for different piezoelectric disks. Data are shown for PZT, PMN-PT ceramic, and PMN-PT single crystal disks [44].

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10.30 V/cm$Oe, which is w80 times higher than that previously reported in the 0-0 particulate composites. This high value of (dE/dH) for PMNPT single crystals is due to the high piezoelectric voltage constant (g33), as well as high elastic compliance (s33). As introduced later in Eq. (10.24), the output voltage from the composite is directly proportional to the piezoelectric voltage constant g31 and the stress on the piezoelectric plate. The stress on the piezoelectric material is dependent on the elastic compliance of the material. The higher the compliance, the higher is the generated stress, as the mechanical coupling is directly correlated with the magnitude of the compliance. Table 10.5 shows the piezoelectric voltage constant and elastic compliance of the three different materials studied in this investigation. It is evident from the table that both the piezoelectric voltage coefficient and elastic compliance are higher for PMN-PT single crystal as compared to PZT ceramic. This difference is responsible for the higher ME properties of the PMN-PT laminate composites. These results indicate that a better ME property can be obtained if the elastic and piezoelectric properties of the materials can be improved.

10.3.4 Thickness ratio effects on the ME properties The ME voltage coefficient was found to increase with decreasing thickness of the PZT layer as depicted in Fig. 10.13(a). This can be explained by the increase in compressive stress in the PZT layer with decreasing thickness of PZT. The compressive stress in the PZT layer and the tensile stress in the Terfenol-D layers can be derived from simple beam theory under plane stress conditions, as indicated in Eqs. (10.22) and (10.23): [47]. sE31t ¼

Et Ep tp Dε0 ; ð1  yÞð2Et tt þ Ep tp Þ

sE31p ¼ 

(10.22)

2Et Ep tt Dε0 ; ð1  yÞð2Et tt þ Ep tp Þ

(10.23)

Table 10.5 Material properties of single and poly piezoelectrics. Material εT33 /ε0 d33 (pC/N) g33 (mVm/N)

sE33 (10L12 m2/N)

PMN-PT single crystal PMN-PT ceramics PZT ceramics

56.4 9.5 17.4

4344 5614 1081

1710 570 250

44.45 11.47 26.11

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(a)

381

(b) 4000

dE/dH (Arb.unit)

3000 2000 1000 0.5 mm PZT (APC840) 0.6 mm PZT (APC840) 0.7 mm PZT (APC840)

0

0

1000

2000

3000

Magnetic field (Oe)

4000

dE/dH

dV/dH

Thick PZT 5000

dV/dH (Arb. unit)

dE/dH (mV/Oe cm)

5000

0

5

Thin PZT

10

15

20

25

Thickness ratio (tt/tp)

Figure 10.13 (a) ME voltage coefficient as a function of applied d.c. magnetic bias with various thickness of PZT layer 1 kHz. (b) Theoretical expectation of the magnetoelectric voltage coefficient as a function of thickness ratio (tt/tp) between Terfenol-D and PZT layer.

where s31 ; Dε0 ; E; t, and y are the transversal stress normal to the PZT disk (i.e., radial direction), linear strain (in-plane) of the Terfenol-D layer, elastic (Young’s) modulus (i.e., stiffness), thickness, and Poisson’s ratio (Poisson’s ratios of Terfenol-D and PZT are assumed to have the same value in these equations), respectively. The subscript t or p stands for Terfenol-D or PZT, respectively. As shown in these equations, the compressive stress in the PZT layer is increased with the decreasing thickness of the PZT layer or the increasing thickness of the Terfenol-D layer. Since the thickness of the Terfenol-D is fixed at 1 mm (in our design), by decreasing the PZT layer thickness, the compressivestress in the PZT layer is increased. The output voltage from the composite can be expressed by the following equations: Vout ¼ 2  g31  tp  sE31p ; 2  g31  sE31p dE Vout ¼ ðV = cm $ OeÞ: ¼ dH Hac  tp Hac

(10.24) (10.25)

Therefore, a higher output voltage can be obtained when the compressive stress in the PZT layer is higher, i.e., a thinner PZT layer. From these equations, it can be seen that the output voltage from the composite is also directly proportional to the piezoelectric voltage constant g31. Generally in PZT ceramics, the piezoelectric voltage constant g31 is around one-third that of g33. In this regard, the laminate composite made with APC-840, which has the highest g33, exhibits the highest ME voltage coefficient, as shown in Fig. 10.11.

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Fig. 10.13(b) shows the theoretical expectation for the ME voltage coefficient (dE/dH) and the output voltage (dV/dH) as a function of the thickness ratio (tt/tp) between Terfenol-D and PZT. The ME voltage coefficient increases with increasing thickness ratio (tt/tp), but output voltage decreases with increasing thickness ratio. Both values saturate above a thickness ratio of 10. The value of the output voltage is more important than the ME voltage coefficient for practical sensor applications. Therefore, a lower thickness ratio (less than 10) is more suitable, even though the ME voltage coefficient is increasing with thickness ratio.

10.3.5 Magnetostriction direction dependence Three kinds of laminate composite design were prepared by using two types of Terfenol-D disks to investigate this issue. These are as follows: (1) Composite with two Terfenol-D disks that have their magnetostriction along the thickness direction (denoted as Comp.T-T). (2) Composite with one Terfenol-D disk with thickness magnetostriction direction and the other Terfenol-D disk with radial magnetostriction direction (denoted as Comp.T-R). (3) Composite with two disks that have their magnetostriction along the radial direction (denoted as Comp.R-R). Fig. 10.14(a) shows schematic illustrations of each composite structure. The dielectric polarization direction of the PZT disk and the applied magnetic field direction were in the thickness direction for all the composites. Fig. 10.14(b) illustrates the variation of the ME voltage coefficient (dE/dH) as a function of the d.c. magnetic bias for the three different composites. The ME voltage coefficients of all composites increased with increasing d.c. bias, and saturated around 4 kOe. The Comp. R-R showed the most superior ME property. Its maximum ME voltage coefficient was 5.90 V/cm$Oe when under a magnetic d.c. bias equal to or greater than 4.2 kOe. In the Terfenol-D disks, the magnitude of strain in the principal magnetostriction direction is higher than other directions [48].

10.3.6 Magnetic field direction dependence In applications like magnetic field sensing devices, the dependence of the ME response on the magnetic field direction is an important factor. To examine the magnetic field direction dependence, we measured the ME voltage coefficient by changing the applied magnetic field direction. The dependence of the ME voltage coefficient of the composites on the applied

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(a) Magnetostriction Terfenol-D direction

Polarization and magnetic field direction

PZT Comp. T-T

(b)

Comp. T-R

7000 6000

dE/dH (mV/cm Oe)

Comp. R-R

Comp T-T Comp T-R Comp R-R

5000 4000 3000 2000 1000 0 0

1000

2000

3000

4000

5000

Magnetic field (Oe)

(c)

3000 Comp T-T Comp T-R Comp R-R

dE/dH (mV/cm Oe)

2500 2000 1500 1000 500 0 0

20

40

60

80

100

Magnetic field direction θ (°)

Figure 10.14 (a) Schematic illustrations for three different PZT/Terfenol-D composites. (b) ME voltage coefficient as a function of applied d.c. magnetic bias with different assembly. (c) ME voltage coefficient as a function of an applied a.c. magnetic field direction.

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a.c. magnetic field direction is shown in Fig. 10.14(c). The angle of the applied a.c. magnetic field (1 kHz) indicates the difference in angle between the a.c. magnetic field and the d.c. magnetic bias (thickness direction of the sample). The dependence of the a.c. magnetic field direction exhibited a similar behavior for all the composites, with a maxima occurring at 25e45 degrees. Beyond 45 degrees, the ME voltage coefficient decreased with an increase of the a.c. magnetic field direction. According to the theoretical calculations, these behaviors are basically related to the areal strain changing behavior of Terfenol-D with changing the applied magnetic field direction. The calculated maximum of the areal strain occurs at the orientation angle, q ¼ 51 degrees. It is expected that the ME coefficient will show a maximum around this angle. This behavior originates from the contribution of the relatively large shear mode strains, i.e., |d15| > |d33| or |d31| [49]. In the end of this Section 10.3, ME laminate composites, the author discloses here the ICAT’s development strategies on the magnetoelectric devices. We initially started to work on various solid solution systems of complex perovskites, which include magnetic ions such as Fe, Co, Ni, Mn, etc., after chasing Smolenskii’s group. Though we could develop “ferrimagnetic-ferroelectric” compositions, the coupling phenomena were not strong, and observed only at low temperature. Because of this discouraging result from the industrial application viewpoints, the research enthusiasm faded out in the early 1980s. We needed to wait for roughly one generation period (i.e., 30 years) until the “research renaissance (revival)” in the early 21st century, now from both industrial (composites for sensors, energy harvesting applications) and scientific (“multi-ferroic” coupling) viewpoints. We now chased the Philips group on the 0-0 type composites with magnetostrictive and ferroelectric/piezoelectric powder mixtures, which exhibited 10 times higher ME voltage coefficients than those of most single-phase materials. Their idea encouraged us to move into the composites. But, even we used PZT, much better piezoelectric property than barium titanate, our device performance exceeded only 40% to the Philips. This discouragement instructed us “not to chase the others idea.” Because the author is the inventor of “multilayer” actuators, and also from the “Connectivity” model analysis/simulation, we concluded to use the laminate composite. Once the research direction was fixed, the approaches were not very difficult. Thanks to our former research associates, various composites were prepared by stacking and then bonding piezoelectric materials to Terfenol-D. Laminates were made using several kinds of PZTs, PMN-PT ceramics, and a

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(001)-oriented PMN-PT single crystal. The highest ME voltage coefficient (dE/dH) was found for the PMN-PT single crystal, which was 10.30 V/ cm$Oe, under a d.c. magnetic bias of 0.4 T, another 10 times higher than the 0-0 type powder mixture composites. This unique laminate design concept created our current engineering leading status. You can easily understand the following systematic optimization processes in the above sections: To obtain excellent ME property from the ME laminate composites, a high piezoelectric voltage coefficient (gij), an optimum thickness ratio between piezoelectric layer and Terfenol-D layers, the direction of magnetostriction in the Terfenol-D disks, and higher elastic compliance of piezoelectric material are important factors. Note also that our laminate composite design fits beautifully with so-called MEMS micromachining technologies for miniaturizing the magnetoelectric devices.

10.4 Magnetoelectric applications 10.4.1 Sensor applications Similar to nuclear radiation, magnetic irradiation cannot be easily felt by humans. Some reports mentioned that brain cancer may be triggered by a frequent usage of a mobile phone, though there is no strong scientific evidence. The problem is the situation where we cannot even purchase a magnetic field detector for a low frequency (50 or 60 Hz). The ME laminate composite developed by the Penn State is for a simple and handy magnetic noise sensor for the environmental monitoring purpose, which was the initial motivation for developing the laminate composites. A Japanese real-estate agency sponsored us the device to monitor the magnetic field below high-power transmission cables, in order to sell houses there without discounting. Refer to Fig. 10.9(b), where we integrated a pair of permanent magnets to provide a suitable bias magnetic field on the ME laminate sensor. Another application is for the Navy; that is, to survey the enemy submarine which equips the acoustic stealth system. Large Fe (magnetic) mass movement (even slow) modulates the Earth’s magnetic field at around 10 Hz or lower frequency, which can be detected by the laminate ME device down to 1 nT level. The key of this device is high effectiveness for a low frequency magnetic field modulation. See Fig. 10.15 for understanding visually. Similar sensor applications include (a) predicting the earthquake

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Don’t worry! magnetic field is low enough !

µ 0I B= 2S r

Magnetoelectric sensor (20u20u100 mm3)

Large Fe (magnetic) mass (with motion, even slow)

Figure 10.15 (a) Stray magnetic field monitoring below high-power transmission cables (50 Hz). (b) Acoustic stealth-submarine monitoring under the sea (w10 Hz or lower).

occurrence through the earth magnetic field fluctuation via magma movement in a weekly level short-term, (b) discovering land-mines and mineral mines, now by sweeping the ME sensor (i.e., crisis technologies).

10.4.2 Energy harvesting applications In addition to monitoring the “stray” magnetic field, electric energy harvesting is also possible from the magnetic field under the power transmission cable, if the field strength is high enough. Fig. 10.16(a) shows the present human manual inspection procedure of the high-power transmission cables. In order to reduce the required manpower, an unmanned vehicle is being introduced by an electric company for checking the cables, as shown in Fig. 10.16(b). Though a rechargeable battery is installed on this unmanned vehicle, it is much more convenient to equip an energy harvesting device from reasonably high magnetic field very close to the power cable.

Figure 10.16 (a) Human manual inspection of high-power transmission cables. (b) Unmanned inspection robot of transmission cables.

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Virginia Tech group proposed a hybrid energy harvesting of the ME type with a piezoelectric type bimorph element. Fig. 10.17 illustrates the cross-section view of the hybrid laminate structure of PZT fiber composite bimorph and Metglas magnetostrictive foils [50]. Under a magnetic field applied, bending vibration of the laminate composite is excited, then the electric power is generated from the piezo-bimorph. Also, under a mechanical vibration, electric voltage is excited from the piezo-bimorph simultaneously. This sort of hybrid energy harvesting seems to be useful for health-monitoring industrial motors in a factory. Since the motor generates both vibration and leak magnetic field adjacent to the motor at the same 50 Hz range, superposed larger electric energy can be harvested for sending the signal of motor condition remotely to the operation room (i.e., machine safety monitoring). Fig. 10.18 summarizes possible hybrid energy harvesting devices based on ferroelectric/piezoelectric plates. In addition to

FeBSiC layer

PZT fiber composite H bimorph

VA

F

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P

Vibration VB

Figure 10.17 ME devices composed with PZT fiber bimorph composite sandwiched by Metglas layers [50].

Piezoelectric plate Elastic plate

• Vibration noise → elastic material → Piezoelectric effect

• Magnetic noise → magnetostrictive material → magnetoelectric effect

• Photo illumination º Pure light → elastic material → photovoltaic effect º Photothermal heat → elastic material → pyroelectric effect

Figure 10.18 Possible hybrid energy harvesting devices based on piezoelectrics.

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magnetoelectric effect, we can couple with various multifunction effects such as photovoltaic and pyroelectric effects to enhance the harvesting energy level.

10.5 Summary 1. Multiferroic coupling can be found between ferroelectric and ferroelastic (piezoelectricity), ferromagnetic and ferroelastic (piezomagnetic, magnetostriction), or ferroelectric and ferromagnetic (magnetoelectric, ME). Strong “direct” ME coupling between ferro- (or ferri-) magnetic and ferroelectric has not been realized at room temperature in a single-phase material. 2. 0-0 connectivity ME composites exhibit 10 times higher ME voltage coefficient at room temperature. The 2-2 connectivity, TerFeNOL: PZT laminates exhibit further significant (80 times higher) ME response, which has triggered the current research boom. 3. Designing principle of ME compositesdoptimization was studied in terms of (1) crystal orientation, (2) layer thickness, (3) DC bias magnetic field, etc. 4. Sensor applications of ME composites include (1) magnetic noise sensor, (2) geomagnetic sensor, (3) earthquake prediction, (4) antistealth submarine technology. 5. ME energy harvesting realizes hybrid energy harvesting systems for stray mechanical vibration and stray magnetic field simultaneously.

Acknowledgments The researches on the magnetoelectric laminate composites are highly indebted to our former associates, Drs. Jungho Ryu (Yeungnam University, Korea) and Shashank Priya (The Pennsylvania State University, USA). We also appreciate for the continuous research support from the US Office of Naval Research via #N00014-17-1-2088.

References [1] J. Valasek, Piezoelectric and allied phenomena in Rochelle salt, Phys. Rev. 15 (6) (1920) 537. [2] K. Aizu, J. Phys. Soc. Jpn. 38 (1975) 1592. [3] P. Curie, J. Physique 3eSeries 3 (1894) 393. [4] I.E. Dzyaloshinskii, Sov. Phys.dJETP 37 (1960) 628. [5] D.N. Astrov, Sov. Phys.dJETP 11 (1960) 708. [6] G.T. Rado, V.J. Folen, Phys. Rev. Lett. 7 (1961) 310. [7] A.J. Freeman, H. Schmid, Magnetoelectric Interaction Phenomena in Crystal, Gordon and Breach Science Publishers, London, UK, 1975.

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[8] E. Ascher, H. Rieder, H. Schmid, H. Stossel, J. Appl. Phys. 37 (1966) 1404. [9] J.P. Rivera, H. Schmid, J.M. Moret, H. Bill, Measurement of the Magnetoelectric Effect in Ni-Cl Boracite”, Chapter Contribution (Page 169) to Magnetoelectric Interaction Phenomena in Crystal, Gordon and Breach Science Publishers, London, UK, 1975. [10] A.F. Devonshire, Phil. Mag. 40 (1949) 1040. [11] A.F. Devonshire, Adv. Phys. 3 (1954) 85. [12] H.F. Kay, Rep. Prog. Phys. 43 (1955) 230. [13] E.F. Bertaut, M. Mercier, Mat. Res. Bull. 6 (1971) 907. [14] R.R. Birss, Symmetry and Magnetism, North-Holland, Amsterdam, 1964. [15] G.A. Smolensky, et al., Segnetoelektriki I Antisegnetoelektriki, Nauka, Leningrad, 1971. [16] E. Ascher, J. Phys. Soc. Jpn. 28 (Suppl. 7) (1970). [17] K. Uchino, S. Nomura, Dielectric and magnetic properties in the solid solution system Pb(Fe2/3W1/3)O3 - Pb(Co1/2W1/2)O3, Ferroelectrics 17 (1978) 505e510. [18] K. Uchino, K. Hoshi, S. Nomura, Dielectric and magnetic properties in the solid solution system Pb(Fe2/3W1/3)O3 - Pb(Mn1/2W1/2)O3, in: Proc. 1st. Meeting on Ferroelectric Mater. Appl, 1977. F11, 321 Kyoto. [19] K. Uchino, S. Nomura, Crystallographic and dielectric properties of Pb(Mg1/2W1/2) O3 e BiFeO3 solid solutions”, Japan, J. Appl. Phys. 17 (8) (1978) 1351e1354. [20] K. Uchino, S. Nomura, Phenomenological theory of ferroelectricity in solid solution systems Pb(Fe2/3W1/3)O3 - Pb(M1/2W1/2)O3 (M ¼ Mn, Co, Ni)”, Japan, J. Appl. Phys. 18 (8) (1979) 1493e1497. [21] K. Uchino, F. Kojima, S. Nomura, Phase transition in the Pb(Fe2/3U1/3)O3 - PbZrO3 system, Ferroelectrics 15 (1977) 69e71. [22] G.A. Samara, J.F. Scott, Solid State Comm. 21 (1977) 167. [23] J.F. Scott, Phys. Rev. B16 (1977) 2329. [24] G.T. Rodo, Phys. Rev. 128 (1962) 2546. [25] E.I. Venturini, F.R. Morgenthaler, in: C.D. Graham Jr., G.H. Lander, J.J. Rhyne (Eds.), AIP Conf. Proc. 24, AIP, New York, 1975 page 168. [26] D.L. Fox, J.F. Scott, J. Phys. C10 (1977) L329. [27] K. Uchino, S. Nomura, “Crystallographic and dielectric properties of Pb(Mg1/2W1/2) O3 - BiFeO3 solid solutions”, Japan, J. Appl. Phys. 17 (8) (1978) 1351e1354. [28] D. Lebeugle, et al., Phys. Rev. Lett. 100 (2008) 227602. [29] Y.F. Popov, et al., JETP Lett. (Engl. Transl.) 57 (1993) 69. [30] A.K. Zvezdin, et al., J. Magn. Magn Mater. 333 (2006) 224. [31] G.T. Rodo, J.M. Ferrari, Phys. Rev. B12 (1977) 5166. [32] N. Shiratori, et al., Proc. Magnetic Section Meeting of Japanese, Phys. Soc., October 1976. [33] Y. Miyamoto, et al., Proc. Magnetic Section Meeting of Japanese, Phys. Soc., October 1976. [34] R.E. Newnham, et al., Mater. Res. Bull. 13 (1978) 525. [35] J. Van Suchetelene, Philips Res. Rep. 27 (1972) 28. [36] K. Uchino, Solid State Phys. 21 (1986) 27. [37] K.A. Klicker, J.V. Biggers, R.E. Newnham, J. Amer. Ceram. Soc. 64 (1981) 5. [38] K. Uchino, Photostrictive effect and its applications, Solid State Phys. 22 (8) (1987) 55e60. [39] J. van den Boomgaard, A.M.J.G. Van Run, J. Van Suchetelene, Ferroelectrics 10 (1976) 295. [40] J. Ryu, A. Vazquez Carazo, K. Uchino, H.E. Kim, Piezoelectric and magnetostrictive properties of lead Zirconate titanate/Ni-ferrite particulate composites, J. Electroceramics 7 (2001) 17e24.

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[41] J. Ryu, A. Vazquez Carazo, K. Uchino, H.E. Kim, Magnetoelectric properties in piezoelectric and magnetostrictive laminate composites, Japan. J. Appl. Phys. 40 (2001) 4948e4951. [42] J. Ryu, S. Priya, A.V. Carazo, K. Uchino, H.-E. Kim, Effect of the magnetostrictive layer on magnetoelectric properties in Pb(Zr,Ti)O3/Terfenol-D laminate composites”, J. Amer. Ceram. Soc. 84 (2001) 2905e2908. [43] J. Ryu, S. Priya, K. Uchino, H.-E. Kim, Magnetoelectric effect in composites of magnetostrictive and piezoelectric materials, J. Electroceramics 8 (2002) 107e119. [44] J. Ryu, S. Priya, K. Uchino, H.-E. Kim, D. Viehland, High magnetoelectric properties in 0.68Pb(Mg1/3Nb2/3)O3-0.32PbTiO3 single crystal and Terfenol-D laminate composites, J. Korean. Ceram. Soc. 39 (2002) 813e817. [45] J. Kuwata, K. Uchino, S. Nomura, Ferroelectrics 37 (1981) 579. [46] J. Kuwata, K. Uchino, S. Nomura, Japan. J. Appl. Phys. 21 (1982) 1298. [47] A.V. Virkar, J.L. Huang, R.A. Cutler, J. Amer. Ceram. Soc. 70 (3) (1987) 164. [48] G. Engdahl, Handbook of Giant Magnetostrictive Materials, Academic Press, San Diego, CA, 2000, p. 127. [49] G. Engdahl, Handbook of Giant Magnetostrictive Materials, Academic Press, San Diego, CA, 2000, p. 175. [50] J. Zhai, P.D. Thesis, Virginia Tech, 2009.

CHAPTER ELEVEN

Magnetic field into multifunctional materials Magnetorheological, magnetostrictive, and magnetocaloric P. Martins1, 2, S. Lanceros-Mendez3, 4 1

Centre/Department of Physics, University of Minho, Braga, Portugal IB-S Institute of Science and Innovation for Sustainability, University of Minho, Braga, Portugal BCMaterials, Basque Center for Materials, Applications and Nanostructures, UPV/EHU Science Park, Leioa, Spain 4 IKERBASQUE, Basque Foundation for Science, Bilbao, Spain 2 3

11.1 Introduction Magnetoactive materials are a group of smart materials that can adaptively modify their physical properties as a response to external magnetic fields, allowing wireless activation, response, and controllability [1]. They are widely used in electric motors, power generators, memory devices, windmills, biomedical devices, energy conversion, and transportation, among others [2,3]. The characteristic fingerprint of magnetism is the existence of an ordered arrangement of magnetic moments over macroscopic length scales, with a spontaneous breaking of symmetry [4]. This is classically driven by the interaction between the neighboring spins that tend to favor specific relative orientations between them. At T z 0 K, this local order can extend over macroscopic length scales. With increasing temperature, thermal fluctuations lead to misalignment of magnetic moments in neighboring regions, so that long-range order disappears above a certain critical temperature (Tc). The possibility of a system to suffer a phase transition at a finite Tc depends, particularly and mostly, on the effectiveness of thermal fluctuations, which are controlled by a small number of general parameters of the system. Specifically, dimensionality issues play a key role in the determination of the impact of thermal fluctuations on the critical behavior of magnetoactive systems [3,4].

Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00009-3

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In a 3D magnetoactive system, the magnetic phase transition can certainly occur at a finite temperature, whereas in a 1D magnetoactive system it is possible only at zero temperature [5]. Between 1D and 3D, the situation in 2D systems is far more complex and intriguing once the existence of magnetic long-range order at any finite temperature critically depends on the number n of relevant spin dimensionality (Fig. 11.1), being determined by physical parameters of the system such as the presence and strength of magnetic anisotropy [4]. In a challenging 4.0 industry paradigm and contextualized in a dynamic Internet of Things (IoT) environment, a new generation of magnetoactive, multifunctional materials, such as magnetorheological, magnetostrictive, and magnetocaloric materials (Table 11.1), capable of changing their physical properties (rheology, strain capabilities, and heat generation, respectively) under the influence of external magnetic fields, are attracting increasing interest due to the wide range of interesting physical phenomena observed in these materials and the large potential of their practical application in technological devices. Magnetorheological materials are a class of magnetoactive smart materials whose rheological properties may be varied by application of a magnetic field. These materials traditionally consist of micron-sized ferrous particles dispersed in a fluid or an elastomer [6]. Thus, the mechanism responsible for this bulk effect is the induced magnetic interaction of particles within the matrix. On the other hand, magnetostrictive materials can be seen as a collection of magnetic domains whose orientations depend on the interplay between magnetic and mechanical energies and whose dimensions may be varied by application of a magnetic field. The magneto-mechanical coupling induces several behaviors that are relevant to structural vibration control: the Joule effect, the Villari effect, material hysteresis, and the Delta-E effect [7]. Ni, Fe, Fe3O4, CoFe2O4, Terfenol-D, Tb0.5Zn0.5, Tb0.5DyxZn are the most studied magnetostrictive materials [8]. It is also important to notice that magnetostrictive materials combined with piezoelectric ones allow the development of magnetoelectric composites already discussed in a previous chapter. Analogously, the magnetocaloric effect is referred as the isothermal change of entropy (DS) and adiabatic change of temperature (DT) upon the variation of the magnetic field. The value of DT is directly measured using a thermometer or indirectly from specific heat data while DS is

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Figure 11.1 Role of spin dimensionality and evolution of Tc. (a) A spin dimensionality n ¼ 1 means that the system has a strong uniaxial anisotropy and the spins point in either of the two possible orientations (up or down) along a given direction. The system behaves effectively as if it has only a single spin component along the easy axis, and the underlying spin Hamiltonian for localized spins is the Ising model. The case n ¼ 2 corresponds to an easy-plane anisotropy that favors the spins to lie in a given plane, although the orientation within the plane is completely unconstrained. The spins can thus be considered to have effectively only two components (associated with the two in-plane directions), which are successfully described within the XY model. Note that in this case magnetic susceptibility (c) / N for T < TKT (KosterlitzeThouless temperature). Finally, for isotropic systems, n ¼ 3 and there is no constraint on the direction of the spins. The underlying spin Hamiltonian in this case is the isotropic Heisenberg model. (b) Tc (normalized to bulk critical temperature 3D for the particular material) as a function of the number of layers. Taken with permission from M. Gibertini, M. Koperski, A.F. Morpurgo, K.S. Novoselov, Magnetic 2D materials and heterostructures, Nat. Nanotechnol. 14 (2019) 408. 393

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Table 11.1 Figures of merit (FOM) of some of the most reported magnetorheological, magnetostrictive, and magnetocaloric materials (cr magnetorheological coefficientdexpressing the variation of the shear stress as a function of the variation of the magnetic field, Sm is the maximum strain in ppm, and RCP is the relative cooling power in J.kg1). Material Type FOM Refs.

ZnFe2O4 Fe3O4 Carbonyl iron/g-Fe2O3 CoFe2O4 MnFe2O4/graphene oxide Ni Fe Fe3O4 CoFe2O4 Terfenol D Tb0.5Zn0.5 Tb0.5DyxZn EuTiO3 HoMnO3 Gd (bulk) Zn0.6Cu0.4Fe2O4 Ni0.5Zn0.5Fe2O4

Magnetorheological

cr ¼ 11 kPa/T cr ¼ 3 kPa/T cr ¼ 34 kPa/T cr ¼ 1 kPa/T cr ¼ 10 kPa/T

[10] [11] [12] [13] [14]

Magnetostrictive

Sm ¼ 50 ppm Sm ¼ 14 ppm Sm ¼ 60 ppm Sm ¼ 208 ppm Sm ¼ 2000 ppm Sm ¼ 5500 ppm Sm ¼ 5000 ppm RCP ¼ 328 J.kg1 RCP ¼ 540 J.kg1 RCP ¼ 690 J.kg1 RCP ¼ 289 J.kg1 RCP ¼ 161 J.kg1

[8] [8] [8] [15] [8] [8] [8] [16] [17] [18] [19] [20]

Magnetocaloric

calculated from magnetization or specific heat [9]. Perovskites (usually structures with the general formula of ABO3), glass composites/alloys, and spinel ferrites are the most studied magnetocaloric materials. In this chapter, a brief history, classification, and state-of-the-art of magnetoactive, multifunctional materials will be presented. Next, it will be discussed the need of the transition from high weight to lightweight and from ideas to applications. Finally, the conclusions and future trends of this dynamical research field will be presented.

11.2 From high weight to lightweight With a unique freedom regarding material selection, magnetoactive composites allow novel designs able to drive operation conditions to new limits, unlocking the potential for novel applications [21]. Although the first idea to apply soft magnetic composites based on iron powders in electrical machinery was produced as early as the 19th century, it did not attract much attention until the 1980s, when Kordecki and

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Weglinski [22] described several soft magnetic powder composites and the problems associated with their potential application as magnetic cores in electrical devices [23]. Since then, and boosted by both the aeronautics and aerospace industries [24], research on the development of lightweight, magnetorheological, magnetostrictive, and magnetocaloric materials and their technological applications has been intensified and strong progress has been achieved. After those first magnetoactive materials reported (namely soft ferrites) good magnetic properties, some shapes and sizes problems emerged, making them unsuitable for effective technological device applications [25]. To solve those problems, polymer-iron-based composites have emerged as a solution, surpassing the problems found in silicon steels and ferrites. The production of these materials usually consists of magnetic powders such as pure Fe or Fe-based alloys with Ni, Co, P, Nd, B, Mn, Zn, and Ba within a polymeric matrix [25,26]. Some studies reported/discussed effects such as compaction pressure [27], sintering temperature [28], annealing treatments [29], polymer content [30] and grain size [31] in order to determine the possibilities and limitation of the developed materials to perform specific functions on device applications [25]. In order to produce highly efficient polymer-based magnetoactive and multifunctional materials, the addition of magnetic nanoparticles made of inorganic matter (traditionally Fe3O4, g-Fe2O3, “soft” magnetic Fe and also “hard” magnetic such as Co, Ni, FeN, FePt, FeP) into polymeric materials has been consolidated as the more appealing and efficient solution [32]. Indeed, this approach allows to produce nano/micro-scaled magnets with magnetic moments higher than those of molecular magnets, allows them to respond to weak magnetic stimulation (static or alternating magnetic field) with a significant effect (change in the rheology, strain or temperature), making them suitable, for example, for drug delivery or separation applications (Fig. 11.2). The combination of the magnetic properties of those fillers with the elastic properties of the polymers leads to prominent new phenomena that are exhibited as a response to external magnetic fields [33]. Other interesting effects such as the giant deformational effects, the high elasticity effect, the anisotropic elastic effect, and the swelling effects linked to a fast response to external magnetic fields and their real-time controllable elastic properties open new opportunities/challenges for using such magnetoactive/multifunctional materials for numerous applications as smart materials in engineering devices [33].

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H dc Controlled drug delivery

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H dc Shape memory polymer materials Molecule separation

Figure 11.2 Schematic representation of different types of magnetic responsive materials obtained from the doping of various polymers with magnetic particles and illustration of their response when exposed to a static magnetic field (HDC) or to an alternating magnetic field (HAC). From left to right: composites made from elastomeric polymers can be deformed in homogeneous fields or gradients in a controlled fashion; magnetoresponsive polymer composite (MRPC) particles made of polymers designed for biomedical applications can be used for magnetic guidance for drug delivery or separation purposes; MRPC from thermoresponsive polymers can be activated by magnetic induction using alternating fields. Reproduced with permission of J. Thévenot, H. Oliveira, O. Sandre, S. Lecommandoux, Magnetic responsive polymer composite materials, Chem. Soc. Rev. 42 (2013) 7099.

More recently, increasing and focused interest has been devoted to development and use of multiresponsive, magnetic, polymer-based composites, which exhibit sensitivity to several external stimuli, namely, magnetic field, temperature variations, and pH changes.

11.3 From idea to applications 11.3.1 Magnetorheological, lightweight materials Magnetorheological, lightweight, multifunctional materials have been used in many applications in recent years. Wilson et al. [34] developed

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polyurethane and silicone polymer gels whose rheology was qualitatively controlled for each system by proper selection of reactants and diluents concentrations. The resulting polymer-based materials exhibited solid, gel, or liquid states, depending on the cross-linking and dilution, with potential applications in vibration control, damping, and energy-absorption. The control of vibrations and switching/control of torque/force in dampers, shock absorbers, isolators, and brakes are some engineering applications that can take advantage of core-shell-structured magnetic carbonyl iron (CI)-poly(methyl methacrylate) (PMMA) particles fabricated via CIseeded dispersion polymerization method, in order to enhance dispersion stability of the magnetorheological fluid when dispersed in mineral oil [35]. A series of magnetorheological, lightweight gels, consisting of plastic polyurethane matrix swollen by nonvolatile solvents in different weight fractions and carbonyl iron particles was prepared [36]. It was discovered that by introducing a gravity yield parameter, the mismatch of density between the carrier medium and the iron particles (essential for particle settling) was improved. Additionally, if the solvent content is lower than 25 wt.% or the gravity yield parameter is higher than 0.865, the polyurethane-based magnetorheological gels will be stable and no particle settling will occur. More important, on the magnetorheological gels with 30 wt.% solvent content, the apparent viscosity can be increased by z 3 orders of magnitude when the magnetic field increases from 0 to 930 mT, showing the strong potential in applications related with smart control. The synthesis and characterization of magnetorheological elastomers with good wettability, good dispersibility, high thermo-oxidative stability, high chemical stability, and sufficient durability were presented by Cvek et al. [37]. This was achieved by adding carbonyl iron particles into a poly(trimethylsilyloxyethyl methacrylate) matrix, using a surfaceinitiated atom transfer radical polymerization technique. An increased particle mobility that has as a consequence an enhanced relative magnetorheological effect (23% higher) was detected. Such fabricated, magnetoresistive, lightweight material containing CI-g-poly(trimethylsilyloxyethyl methacrylate) exhibiting multifunctionality (wettability, magnetorheology, and magnetostriction) can find applications both in damping systems (high thermo-oxidation stability and antiacid/corrosion properties) and in magnetostriction-based sensors (Fig. 11.3). Magnetorheological, lightweight materials were also recently used as a new tool for controlling the performance of oil reservoirs and reducing hazardous water production. This was achieved by using polyacrylamide-coated

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covalent bonds Bare CI

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4u10 3u10 2u10 1u10 0

H 0

200 400 600 800 Magnetic field (kA/m)

Relative MR effect Bare CI – 42.8% CI-g-PHEMATMS – 66.0%

Figure 11.3 Schematic representation of the CI-g-poly(trimethylsilyloxyethyl methacrylate) (PHEMATMS) magnetoactive, multifunctional materials: wettability, magnetorheology, and magnetostriction. Adapted with permission from M. Cvek, M. Mrlík, M. , J. Mosna 
cek, L. M€ Il
cíkova unster, V. Pavlínek, Synthesis of silicone elastomers containing silyl-based polymer-grafted carbonyl iron particles: an efficient way to improve magnetorheological, damping, and sensing performances, Macromolecules 50 (2017) 2189.

magnetite nanoparticles synthesized using a facile one-step method. The reported strong, magnetorheological responses highlighted the feasibility of a conformance control fluid for making a solidlike structure to block the high permeable zone in oil reservoirs [38].

11.3.2 Magnetostrictive, lightweight materials The most cited paper regarding the application of magnetostrictive, lightweight materials was published in 1999 by Lim et al. [39]. A high value of magnetostriction (536 ppm) at a magnetic field of 1.1 kOe, combined with a high maximum slope in the magnetostriction-applied magnetic field curve of 1.3 ppm/Oe (16.3 nm/A) was verified for the Terfenol-D composite fabricated with an average particle size of 137.5 mm, a phenol content of 3.1 wt.%, and a compaction pressure of 0.5 GPa. At the same time amorphous ribbons of an alloy with the composition (Tb0.33Fe0.67)0.98B0.02, which exhibits good magnetostrictive properties at low magnetic fields, were bonded with a phenol-based binder to fabricate bulk composites [40]. A magnetostriction of 493 ppm (at a magnetic field of 1.1 kOe) was achieved on the lightweight composites, together with a high dl/dH (magnetostriction sensitivity with the applied magnetic field) of 1 ppm.Oe1. As in the previous case of magnetorheological materials, gels have been also produced in order to develop magnetostrictive, lightweight materials [41]. Such materials were based on silicon gel with 80 wt.% of 3.8 mm embedded carbonyl iron particles. Under a magnetic field of 0.44 T the

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compression modulus reaches a value of 861 kPa, which represents a 100% increase from the case at B ¼ 0 T. Such properties can offer potential applications as variable stiffness components, large strain actuators, electromagnetically active damping elements, and artificial muscles [41]. In the context of wearable IoT devices, magnetostrictive wire/polymer composites were developed by embedding Fe-Co wires in an epoxy matrix (Fig. 11.4), and their inverse magnetostrictive characteristics were evaluated. It was reported that the output voltage (z10 mV) of this composite in the compression mode strongly increased with increasing stress rate, opening new application perspectives for the development of lightweight, robust, and efficient magnetoactive energy harvesting devices [42]. Spherical, granular, and flake-type powders were also used to produce polymer-bonded composites with a Fe-Co based alloy, being observed that their magnetostrictive sensitivity at low applied fields was optimized for the composites fabricated with flake-type powders [43]. Additionally, a high magnetostriction of 200 ppm was obtained under an applied magnetic field of 4.3 kOe [43]. Other applications of magnetostrictive, lightweight, multifunctional materials include sensing of ultrasonic waves, actuators, biomedical materials, sensors for magnetic fields and electrical currents, vibration isolation and active control, magnetoelectric sensors and devices, stress sensing, and health monitoring [44].

Cure agent Fe29Co71 Wire

Lower demagnetization effect Stronger magnetocrystalline anisotropy Jig

Mixture

Curing 23°C / 24 h

Prestress design

Epoxy resin

Post-curing 80°C / 3 h

Further adhesive property increase

Tensile load Fe29Co71 wire by drawing

Figure 11.4 Schematic representation of the sample manufacturing procedure of FeCo wires/epoxy composites. Adapted with permission from F. Narita, Inverse magnetostrictive effect in Fe29Co71 wire/polymer composites, Adv. Eng. Mater. 19 (2017).

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11.3.3 Magnetocaloric, lightweight materials

Dispersio n polymeriz atio

n

S

Se

lf-

as

se

m

bl

y

Most of the applications developed from magnetocaloric lightweight materials have been developed over the past decade. A good example was reported by Fujita et al. [45] in 2009, showing the control of the magnetocaloric effect by partial substitution in itinerant-electron metamagnetic La(FexSi1-x)13 for applications in magnetic refrigeration. Li et al. [46] explored the concepts of magnetocaloric effect and lower critical solution temperatures (LCST) in order to develop magnetothermally responsive nanocarrier for magnetothermal drug release under alternating magnetic field (Fig. 11.5). For that, Mn0.2Zn0.8Fe2O4 nanoparticles with low Curie temperature (TC) were dispersed in a polymeric matrix consisting of N-isopropyl acrylamide (NIPAAm) and N-hydroxymethyl acrylamide (HMAAm). A maximum self-heating temperature of 42.9 C was achieved by optimizing the nanoparticle content (8wt.%) in the polymer matrix. Additionally, a good biocompatibility and efficient therapeutic effects in cancer treatments were observed, presenting high potential in clinical systemic therapeutics.

N

Magnetocaloric effect Drug release

T > LCST Polymer shrink

Figure 11.5 Schematic illustration showing the preparation process of the nanocarriers and mechanism of magnetothermal drug release. Adapted with permission from J. Li, Y. Qu, J. Ren, W. Yuan, D. Shi, Magnetocaloric effect in magnetothermally-responsive nanocarriers for hyperthermia-triggered drug release, Nanotechnology 23 (2012).

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An innovative mode of radical polymerization (by the assembly of “host” b-cyclodextrin (CD) and “guest” N-vinylimidazole (VI)) was demonstrated by Yu et al. [47], which was accomplished through the magnetocaloric effect, aiming to fast fabricate (5 min) novel hosteguest supramolecular gels. The resulting Fe3O4-doped gels were found to have self-healing properties spontaneously induced when damaged and under an applied magnetic field (450 kHz). Such strategy might open a promising way for accelerating the use of hosteguest assemblies to quickly build bio-inspired robust materials. Radulov et al. [48] recommended the use a particle sizes larger than 200 mm, a compaction pressure of 0.1 GPa, and approx. 5 wt.% of low viscosity epoxy adhesive (silver-based epoxy H27D and epoxydharz L) La (Fe, Mn, Si)13Hx composite for optimized magnetocaloric properties (DT ¼ 4.8 K for a Dm0H ¼ 1.9 T), suitable for magnetic refrigeration at room temperature. Multifunctional hydrogels with both vivid color change and shrinkingswelling response (with clear and fully reversible colorimetric sensing ability in the temperature range 15e32 C and NaCl concentration from 200 to 1200  103 M (ion strength), respectively) to external stimulus such as temperature, ion strength, and alternating magnetic field were produced through magnetic assembly [49], with potential applications such as target medical therapy, tissue engineering, wound healing, and in vivo drugdelivery studies. Still in hydrogels, Yu et al. [50] demonstrated the fabrication of quadruple stimuli-responsive hydrogels with self-repair capacity via rapid interface-directed frontal polymerization. The as-prepared hydrogels, composed of 2-hydroxypropyl acrylate (HPA), 1-vinyl-2-pyrrolidinone (NVP), and graphene oxide, revealed auto-healing without the assistance of any external stimuli and the addition of graphene oxide can lead to better performance in toughness and healing efficiency (91.5%), being the magnetocaloric effect used to contribute to the ignition process in the oil phase. Epoxy (Amerlock Sealer)-bonded La-Fe-Co-Si magnetocaloric plates were presented by Pulko et al. [51]. The resulting assembly fabricated in the form of stacked epoxy-bonded plates was fabricated and tested on an experimental magnetic cooling device, being reported a maximum temperature span of approximately DT ¼ 10 K under magnetic field change of m0H ¼ 1.15 T and a long-term cyclic loading, which is essential for its application in device applications such as magnetic refrigerators.

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Polymer-based (our two-dimensional heterometallic Cu-Ln coordination polymers based on 2-methylenesuccinic acid (H2MSA) ligand, {[Ln2Cu(MSA)4(H2O)6]$2H2O}n(Ln ¼ La (1); Gd (2); Tb (3); Dy (4))) composites were also developed by Li et al. [52], in order to develop magnetic refrigerants for low-temperature applications. The existence of weak magnetic interactions (antiferromagnetic Gd/Gd and ferromagnetic Cu/Gd couplings) and a large magnetocaloric effect (DSm ¼ 36.05(1) J/K kg1) for a DH ¼ 7T is reported.

11.4 Final remarks and future perspectives In the last decade, the use of lightweight, magnetoactive materials has grown both in volume and diversity. The materials development must meet challenging and demanding requirements. This has led to an enhancement in the understanding of the characteristics which are responsible for the individual material magnetoactive properties. In this scenario, both basic knowledge and device integration and applications have stimulated each other so that a large volume of scientific and technological work has been performed [53]. Experimental research must explore a wider range of compounds, taking advantage of a high number of candidate lightweight materials available, with a clear predominance of polymer-based magnetoactive materials. This strategy would allow to structure the magnetic knowledge in magnetoactive materials, leading to fresh insights of the microscopic behavior that should be determined by the strength of the exchange interactions and uniaxial anisotropy [4]. Additionally, future experiments will require considerable technical development on the fillers optimization and new composite ideas to probe magnetism quantitatively on very small length scales and in the conditions of very large demagnetization fields [4]. The magnetic field into multifunctional, lightweight materials considerations taken in this chapter reveals that there are many different directions for future works. The promising results reported over the past years represent the starting point of a new field in which major applications/developments are expected. At this stage, the main questions that are being discussed are essentially of fundamental nature, but as soon as lightweight, magnetoactive materials can be reliably produced with the use of additive manufacturing techniques, based on environmental friendly particles, polymers and solvents, as well as with low particle-to-composite magnetic energy loss, applications based on those materials will be a reality.

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Acknowledgments The authors thank the FCT- Fundaç~ao para a Ciência e Tecnologia-for financial support in the framework of the Strategic Funding UID/FIS/04650/2019 and under project PTDC/ BTM-MAT/28237/2017 and PTDC/EMD-EMD/28159/2017. P. Martins thanks FCT for the contract under the Stimulus of Scientific Employment, Individual Support e 2017 Call (CEECIND/03975/2017). The authors acknowledge funding by the Spanish State Research Agency (AEI) and the European Regional Development Fund (ERFD) through the project PID2019-106099RB-C43/AEI/10.13039/501100011033 and from the Basque Government Industry and Education Department under the ELKARTEK, HAZITEK and PIBA (PIBA-2018-06) programs, respectively. Funding from the European Union’s Horizon 2020 Program for Research, ICT-02-2018 - Grant agreement no. 824339 e WEARPLEX is also acknowledged.

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CHAPTER TWELVE

Multifunctional materials based on smart hydrogels for biomedical and 4D applications 1, 2
Sheila Maiz-Fern
andez1, 2, Leyre Pérez-Alvarez , 1, 2 1, 2 Leire Ruiz-Rubio , José L. Vilas-Vilela , Senentxu Lanceros-Méndez1, 3 1

Basque Centre for Materials, Applications and Nanostructures (BCMaterials), Leioa, Biscay, Spain Department of Physical Chemistry, Faculty of Science and Technology, University of Basque Country, Leioa, Biscay, Spain 3 IKERBASQUE, Basque Foundation for Science, Bilbao, Spain 2

12.1 Introduction The human body presents complex functions, tissues, and organs that are difficult to imitate, replace, or regulate with conventional materials. For this reason, the development of biofunctional materials able to interact with the surroundings, control active agents release, self-repair, dynamically avoid or correct the failure; nowadays continues being a challenging issue. Remarkably, a great effort is currently undertaken for the search of materials that combine with cells to provide differentiated tissue [1]. Tissue engineering is the interdisciplinary field that applies the principles of materials science and biosciences toward the development of technologies that can restore, maintain, and improve partially or totally cellular tissues. These technologies of regenerative medicine have the potential to improve or replace biological functions; hence the great interest that currently arises from their development. The fundamental component of most tissue engineering strategies is the creation of a cellular scaffold; that is, three-dimensional (3D) architectures consisting of structural materials, usually polymeric and biological materials, such as cells, proteins, or growth factors, and in some cases even functional materials such as conductive or magnetic particles. Ideally, the properties of these scaffolds should resemble those of the extracellular matrix and should be designed to initially contain cells and release bioactive molecules [2].

Advanced Lightweight Multifunctional Materials ISBN: 978-0-12-818501-8 https://doi.org/10.1016/B978-0-12-818501-8.00010-X

© 2021 Elsevier Ltd. All rights reserved.

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Thus, the proliferation of investigations on a wide range of natural and synthetic polymers for biomedical and tissue engineering applications is noteworthy. Many of these polymers are materials that respond to chemical or physical stimuli. They are defined as stimuli-responsive polymers or smart polymers. Among this class of polymers, stimuli-responsive polymeric hydrogels have received great attention due to their unique soft nature and water affinity, which make them excellent materials for biomedicine. Polymer hydrogels are hydrophilic three-dimensional cross-linked networks with a high capacity to retain large amount of water, which gives them a characteristic similarity with natural tissue [3]. This ability to absorb liquids is a consequence of the presence of hydrophilic groups (amino, carboxyl, hydroxyl groups, sulfate .) on polymer chains, especially those able to ionize in the presence of water, leading to electrostatic repulsions that result in the swelling of the network. On the other hand, physical or chemical cross-linking allow maintaining the original shape and enable modulating key physical properties of the hydrogels such as swelling, mechanical strength, or elasticity [3e5]. Hydrogels were in the early 1960s one of the first biomaterials specifically synthesized for clinical use [6], and along the next decades, they found interesting applications as wound healing [7], contact lenses [7], superabsorbents [8], tissue engineering [9], cell immobilization [10] and drug delivery systems [11]. Hydrogels have also attracted great interest in other fields such as industry or cosmetics, metals and pollutant recovery [12], or dermal fillers [13]. In addition to the exceptional properties derived from water affinity, smart or stimulus-responsive hydrogels present an additional capability to answer to external changes in environmental conditions. Smart hydrogels can undergo a reversible volume phase transition activated by external physical stimuli, including temperature [14,15], electrical [16], or magnetic fields (Yuhui Li et al., 2013); light or pressure [17] chemical stimuli, such as, pH, ions or ionic strength [18,19]; and biological stimuli such as specific enzymes or biological molecules [20]. The response to a stimulus is mainly controlled by polymeresolvent and/or polymerepolymer interactions which are governed by the chemical structure of the network. A pioneering work on smart hydrogels investigated for their responsive properties corresponded to hydrogels of acrylic acid derivatives that could undergo swelling changes according to the external pH [21]. Since then, more complex hydrogels have been designed to respond simultaneously even to more than one stimulus [22].

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Hydrogels may be undegradable, and consequently stable or permanent, or they can be degraded in the physiological medium; typically these are the cases or synthetic and natural polymers. Nevertheless, synthetic polymers have been specifically modified in order to introduce hydrolysable segments that allow modulating hydrogel disintegration under physiological conditions. Natural hydrogels are typically composed of natural polymers, such as polysaccharides (collagen, hyaluronic acid, alginate, or chitosan) [12] or polypeptides [23]. More used synthetic polymers are those derived from acrylic/methacrylic acid [24], N-isopropylacrylamide [25], and polyethylene glycol [9]. As has been commented, hydrogels display unique characteristics in comparison with other types of biomaterials; however, due to their fragile nature derived from poor mechanical properties, their applicability can be restricted in some cases. Consequently, novel biocompatible hydrogels with tunable biodegradability and improved mechanical and functional properties are synthesized by a combination of naturally formed hydrogels and synthetic hydrogels. Hydrogels are classified as reversible, also called physical gels, when polymeric chains are held together by physical interactions. This is the case of macromolecular entanglements, electrostatic interactions, H-bonding, or hydrophobic forces. A well-known example is the hydrogel formed by the electrostatic interaction between calcium cations and carboxylate anions of ionized calcium alginate. It is said that physical hydrogels usually are reversible because physical interactions can be disrupted by changing the conditions that make them, such as pH or ionic strength (this is the case of calcium alginate) or temperature. Meanwhile, hydrogels are called permanent or chemical gels when the cross-linking of the network is generated by covalent unions between the polymeric chains. These covalent networks can be prepared by the direct polymerization of the monomers in the presence of a multifunctional monomer or cross-linker agent, a well-known example is 2-hydroxyethyl methacrylate (HEMA) copolymers cross-linked with ethylene glycol dimethacrylate (EGDMA) and its derivatives [26], or by cross-linking of polymers by chemical reaction, as for instance in the case of hyaluronic acid with 1,4butanediol diglycidyl ether (BDDE) [27e29]. In some exceptional cases cross-linking is not required. For instance, physical gelation of poly(acrylonitrile) (PAN) takes place after its hydrolysis in which nitrile groups are partially substituted by amide/carboxylic acid groups leading to a specific

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hydrophilic/hydrophobic balance that leads to water retention in a network stabilized by hydrophobic interactions [30]. Regarding the macromolecular structure, different possibilities are found either for physical or chemical hydrogels. Hydrogels can be formed by linear or branched homopolymers or copolymers that can be blended or even different cross-linking can be combined at the same time forming interpenetrating networks (IPNs) [26,31]. As the use of hydrogels for biomedical applications expands, the desired properties of the hydrogel become more specific and complex. In this regard, stimuli-responsive hydrogels are very suitable biomaterials to mimic dynamic aspects of environments in vivo, due to their high water content, elasticity, their ability to store and release biochemical content, and their three-dimensional structure. Indeed, manufacturing using hydrogels as bioinks for 3D printing is a great revolution, since it provides an interesting opportunity to prepare hydrogels with complex structures for tissue engineering and regenerative medicine [7]. However, the systems obtained by 3D printing of conventional hydrogels may not satisfy the need for use in biomedicine. Fortunately, 3D printing also allows the fabrication of biomimetic scaffolds composed of hydrogels that can also dynamically respond by shape changes induced by external stimuli in the microenvironment. This is the recently developed four-dimensional (4D) printing technology consisting of 3D printing of objects able to respond to variations in environmental conditions that trigger the response through conformational changes, deformations, or volume/phase transitions, changing their shape from a temporal shape to a permanent shape. 4D hydrogel arouses great expectations in the form of dynamic substrates capable of effectively mimicking native biological microenvironments [32]. The driven force for the temporal shape transformation of 3D-printed hydrogels for the development of 4D hydrogels is basically based on their intrinsic swelling. In this sense, one of the simplest mechanisms consists of the construction of layered substrates formed by the deposition of materials with differential swelling properties. For instance, Jamal et al. [33] were one of the first in exploiting the differences on the swelling of multilayered hydrogels to get the self-folding in the micrometric scale with predictable curvatures in water solutions. In this example hydrogels were formed by photo-cross-linking of polyethylene glycol (PEG) with different molecular weights that led to variable swelling of bilayered hydrogels, and consequently, to the controlled self-holding of the substrate. Although

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publications on 3D hydrogels have proliferated in the recent years, little has been explored in the use of 3D printing in the creation of 4D hydrogels with acceptable mechanical properties [34]. It is necessary that hydrogels meet the requirements to be adapted for current 3D printing in which the ink is dispensed through a nozzle, by microextrusion or inkjet methodology, and it is deposited on a surface [35,36]. Thus, the election of the hydrogel employed as bioink is a limiting aspect, and for this a wide range of physical, chemical, and biological characteristics have to be considered. These characteristics include required gelation time and gelation mechanism, swelling capacity, degradation rate, viscosity, shear-thinning, viscoelasticity, and biocompatibility [37]. Taking all this into account, progress in current biomedicine and tissue engineering requires a deeper insight on biocompatible stimuli-responsive hydrogels with high cytocompatibility and tunable mechanical properties able to induce programmed phase/shape change under physiological stimuli. In this line, the present work aims to summarize and revise recent investigations on hydrogels sensitive to temperature, light, pH, salt, and biological molecules, in order to offer a broad vision of their preparation methods, properties, and applicability as 4D systems.

12.2 Thermoresponsive hydrogels Some polymers are capable of undergoing solubility changes and phase transitions in response to temperature. In this field, that temperature is known as upper or lower critical solution temperature (UCST and LCST, respectively) [14,15]. In general thermodynamic terms, UCST would refer to the maximum temperature at which the separation of two components occurs, being above it where they would be completely miscible. This is mainly due to the fact that the greater thermal movement exceeds the potential energy that tends to keep molecules of the same type together, while the LCST corresponds to the inverse process, which is exploited to originate induced gelation processes, by changing temperature from room temperature to body temperature [38]. More specifically, in polymers, this phase transition is attributed to a change in the balance between the interactions of the hydrophilic and hydrophobic components of the macromolecule with respect to their interaction with the solvent. In the case of the LCST in polymers, the interactions by hydrogen bonds between the polymer and the water predominate over the hydrophobic interactions at temperatures lower than the LCST, favoring the solubility of

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the polymer. On the other hand, when the LCST is exceeded, the hydrogen bonds weaken and the hydrophobic interactions between polymer chains predominate [39,40]. This response to thermal environment is the socalled hydrophobic effect, and it results in reversible hydrogels which can therefore return in solution after environmental temperature changes [41,42]. This is very useful because temperature is the only stimulus which takes part in their gelation with no other requirement for chemical or environmental treatment, and as a consequence, by an increasing of temperature from ambient to physiological temperature, gels can be formed by injection to the body [15]. On the other hand, thermoresponsive behavior not only depends on solvent interaction with the polymer and hydrophobic/hydrophilic balance with polymer molecules, but also on polymer/solvent interactions which can be varied by adding salts, surfactants, and cosolvents. Among them, surfactants have particular interest, because as soon as surfactants are added to a polymer solution, they have the ability to change the hydrophobic/hydrophilic balance due to their amphiphilic character. The addition of surfactants also can lead to micellization aggregation forms in contrast to a coil to helix transition [43]. One of the most extensively researched thermosensitive polymers is poly(N-isopropylacrylamide) (pNIPAAm), since it is a nonbiodegradable synthetic polymer that undergoes a phase transition at temperatures close to physiological temperature, around 32 C (Fig. 12.1). This phase transition occurs mainly due to the hydrophobic interactions between the isopropyl groups of the polymer chains above the LCST, which leads to the insolubility of the material and, therefore, to its solidification [40]. This synthetic

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Figure 12.1 Coil-to-globule phase transition of PNIPAAm.

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polymer has been used for the manufacture of injectable hydrogels by copolymerization, generally, with natural biopolymers in order to achieve biodegradability property [15,44]. This is the case reported by Ha et al. [45] who prepared a copolymer of PNIPAAm and hyaluronic acid to form thermosensitive injectable hydrogels. For this, thermosensitive polymer was functionalized with terminal amine groups, using 2-aminoethanethiol hydrochloride as a chain transfer agent, which subsequently formed amide bonds with previously activated carboxylic groups of hyaluronic acid. These authors developed a thermosensitive gel with LCST temperature around 33 C, which also possessed good biodegradability profile for drug delivery applications. PNIPAAm is the most prominent candidate as thermoresponsive polymer, even though a second polymer in this class of poly(N-acrylamides) has a nearly identical transition temperature: poly(N,N-diethylacrylamide) (PDEAAm). However, the transition temperature of PDEAAm depends on the tacticity of the polymer, unlike PNIPAAm, which in addition, presents better biocompatibility profile [46] and therefore, its use prevails over PDEAAm for biomedical applications [47]. On the other hand, synthetic amphiphilic polymers have also been used for the development of thermosensitive hydrogels. One of the typical thermosensitive systems are the poloxamers, concretely Poloxamer 407, also known as Pluronic. This type of active agents are ABA type triblock copolymers; they are composed of poly(ethylene oxide) (PEO) (A) and poly(propylene oxide) (PPO) (B) monomers and their aqueous solution can create thermosensitive hydrogels. PPO polymer chain is relatively more hydrophobic than PEO chains and, therefore, when the copolymer is added to an aqueous medium at a concentration above the critical micelle concentration (CMC), it aggregates to form self-assembled micellar structures [48]. These micelles undergo sol-gel phase transition in response to the temperature, which can be modulated by varying the molecular weight of the polymer, its composition, and concentration [15]. An increase in temperature causes the CMC to decrease and, therefore, at moderately low temperatures (T < Tsol-gel) the formation of micelles occurs while at higher temperatures (T > Tsol-gel) the packing of the same originates macroscopic gels (Fig. 12.2) [39,49]. However, like PNIPAAm, the biomedical applications of Pluronic are also strongly limited by their lack of biodegradability and, therefore, it is generally copolymerized with biodegradable blocks, such as, poly(caprolactone), poly(hydroxybutyrate), or poly(acrylic acid) (PAA) [50e52]. The first Pluronic-poly(acrylic acid) research was published by Hoffman et al. [53] In this investigation poloxamer grafted with PAA presents a synthesis method

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Figure 12.2 Sol-gel transition of Pluronic hydrogels.

based on three stages with some intermediate steps, which make it an unattractive method to produce this copolymer on an industrial scale [54]. Nonetheless, taking this fact into account, other researchers like Bromberg et al. [55] developed single-step synthesis method for Pluronic-g-PAA copolymer which was reproducible and scalable. Thanks to Bromberg et al. [55] detailed research, it is well known that poloxamer-g-PAA hydrogels have a thermosensitive sol-gel transition and gel strengths which are dependent on polymer concentration, that is, by increasing the amount of polymer it can get lower transition temperature and higher gel strength [40]. The study of the biocompatibility of hydrogels based on Pluronic has aroused great interest in the biomedical field. Some decades ago, Schmolka et al. [56] provided the first toxicity data of Pluronic which indicated that this triblock polymer was well tolerated. Therefore, this kind of hydrogels were further investigated in medical, pharmaceutical, and cosmetic areas for the treatment of burns [56], for topical administration of anticancer agents [57] or controlled drug delivery, for instance Refs. [57]. The controlled release of this polymer never exceeds a few days according to some publications, and owing to this, Pluronic hydrogels are commonly used for short-term therapies like infections [58] or pain management [59]. Natural polymers have also been studied in recent years to be used as materials for the development of thermosensitive hydrogels. One of the investigations that has caused the most impact in this area is the one published by Chenite et al. [4,60] in which they develop a thermosensitive injectable matrix based on chitosan. Chitosan is a linear polysaccharide, which has ionizable amino groups, so that at pH values lower than its pKa (6.5) [61] it is a polycation, while at pH values above its pKa it is able to create inter/intramolecular hydrogen bonds, which make chitosan insoluble in most solvents and conditions, except in acidic solutions [5,62]. Under these conditions the amino groups are protonated (-NH3þ) and electrostatic repulsions are created along polymer chain, which causes the attractive hydrogen bonds between

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polymer chains to break down, which makes it possible to keep chitosan in solution. Chenite et al. [4,60] solved this problem by adding to the polysaccharide acid solution a basic salt of polyol (b-glycerol phosphate disodium salt). The particularity of this hydrogel is that the addition of the basic salt provides the ideal buffering and control of the hydrophobic interactions and hydrogen bonds, which are necessary to keep the chitosan in solution at neutral pH, since, the hydrogen bonds created with the polyol salt disfavor the intermolecular chitosan-chitosan forces at relatively low temperatures. However, an increase in temperature favors the chitosan-chitosan interaction, and it results in a spontaneous gelation that starts around body temperature (Fig. 12.3). They measured variations in the elastic modulus from (a)

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Figure 12.3 Sol-gel transition of chitosan/b-glycerol phosphate disodium salt hydrogel. (Reproduced with permission from S. Supper, N. Anton, N. Seidel, M. Riemenschnitter, C. Schoch, T. Vandamme, Rheological study of chitosan/polyol-phosphate systems: influence of the polyol part on the thermo-induced gelation mechanism, Langmuir 29 (32) (2013) 10229e10237, https://doi.org/10.1021/la401993q. Copyright © (2013) American Chemical Society).

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0 Pa at 35 C to 1000 Pa at 40 C, and observed that the obtained gels had good mechanical properties and that the gelation temperature could be tuned at 37 C in a short period of time (7 days) and the studies of cytotoxicity and biocompatibility demonstrated the possibility of embedding living cells in the same optimized printing conditions. As it has been seen, thermosensitive hydrogels are based on reversible physical interactions and due to this property these hydrogels offer low stability and poor mechanical properties, which are not able to mimic body tissues. For this reason many authors have chosen to fabricate hydrogels with dual response, that is, to combine sensitivity to temperature with sensitivity

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to other stimuli, for example, light. In this field, there have been reported many works, for example, the one which studied the thermal and photopolymerizable matrix based on poly(N-2-hydroxypropyl methacrylamide lactate) and poly(ethylene glycol). This hydrogel was obtained by thermal gelation and it was endorsed with additional chemical cross-linking by photopolymerization by methacrylate moieties. Moreover, the resulting hydrogels had reproducible porosity, which is an important point to maintain a constant separation between cells or other biological agents. These hydrogels also provided a slow degradation profile and chondrocytes cell culture demonstrated that this polymeric matrix is an interesting candidate for bioprinting [79]. One of the researches that has revolutionized the area of thermoresponsive materials for biomedical applications is the research of 4D bioprintable materials. One of the investigations that has caused the most impact in the field of 4D bioprinting is the one published by Lei et al. [34] in which they attempted to synthesize soft th`ermosensitive sensors that could imitate human skin, that is, appropriate sensory capabilities and mechanical properties. These authors carried out this research based on the problems generated in other previous investigations, such as, poor mechanical stability to mimic tissue and very limited sensory capability. Owing to this, Lei et al. [34] proposed to incorporate a thermosensitive and bioprintable material into a multifunctional system, which is based on calcium carbonate nanoparticles physically cross-linked with poly(acrylic acid) and alginate chains that gave to the material the possibility to sense temperature and body movement to provide sol-gel transition. As said before, the transformation principle of thermoresponsive materials is based on their wettability and solubility changes with the temperature change. Yi Chen-li et al. fabricated double-cross-linked hydrogel based on two components with dramatically different thermal properties induced by crystallization of one of each component [80]. According to this concept, these authors fabricated a hydrogel based on poly(vinyl alcohol) (PVA), which is able to crystallize with chemically cross-linked PEG [50e52]. They obtained stabilized helix morphology after three cycles of freezing when the hydrogel was composed by 70% of PVA and 30% of PEG. Then, the hydrogel was totally capable of transforming this helicoidal morphology to a straight line after 15 s of immersion in water at 90 C, in which the crystalline domains of PVA melted (Fig. 12.4). In addition, deformation angle and diffraction of X-ray revealed that crystallinity could increase when freezing cycles increased and because of this, PEG only with

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Figure 12.4 (a) Photographs of thermal transformation of poly(ethylene glycol) (b) Graph of deformed angle of the hydrogels vs. temperature and (c) Different X-ray diffraction patterns of poly(ethylene glycol) under different freezing/thawing cycles. (Reproduced with permission from G. Li, H. Zhang, D. Fortin, H. Xia, Y. Zhao, Poly(vinyl alcohol)-poly(ethylene glycol) double-network hydrogel: a general approach to shape memory and self-healing functionalities, Langmuir 31(42) (2015a) 11709e11716; J. Li, M. Stachowski, Z. Zhang, Application of responsive polymers in implantable medical devices and biosensors, in: Switchable and Responsive Surfaces and Materials for Biomedical Applications, 2015b, pp. 259e298; L. Li, Y. Wang, L. Pan, Y. Shi, W. Cheng, Y. Shi, G. Yu, A nanostructured conductive hydrogels-based biosensor platform for human metabolite detection, Nano-Micro Lett. 15 (2015c) 1146e1151, https://doi.org/10.1021/nl504217p. Copyright © (2015) American chemical society).

a cycle showed fast response in comparison with that having experimented more cycles [81]. These results confirm that thermoresponsiveness degree of hydrogels could be modulated taking into account the crystallinity of the polymers that are forming the hydrogel. On the other hand, as it is mentioned above, pNIPAAm is known as one of the most important polymers in the field of thermoresponsive materials [82]. Because of this, in pNIPAAm hydrogel polymer chains became hydrophilic when environmental temperature is below LCST. Breger et al. [83] photo-cross-linked acrylic acid functionalized with pNIPAAm (pNIPAAm-AAc) to polypropylene fumarate (PPF). Thus, when the temperature was above 36 C, pNIPAAm-AAc component became to a hydrophobic state to expose the PPF segments so that the shape transformation takes place. At the same time, water was also expelled from the hydrogel as the temperature was increasing; as a result, the change in swelling properties leads to a shape change as can be seen in Fig. 12.4 [84]. In the same area of self-folding systems, Stoychev et al. [85] have designed a star-shape biodegradable and thermoresponsive self-folding bilayered hydrogel based on polycaprolactone (PCL) and poly(N-isopropylacrylamide) which are able to release bioactive materials, such as, cells. Swelling/deswelling

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mechanism depends on pNIPAAm uptake, because of the thermoresponsive properties of this polymer, which undergoes sol-gel transition near physiological temperature. On the other hand, PCL layer acts against swelling in one direction. Therefore, star-shape hydrogel does not present a uniform swelling but it folds/unfolds as a result of collapsing pNIPAAm layer (Fig. 12.5).

12.3 pH responsive hydrogels Among the stimuli-responsive hydrogels, those that present pH sensitivity have been widely studied due to their great applicability mainly as biomaterials. Several pH variations occur in the body, the pH of saliva being (6.5e7.5), upper stomach (4.0e6.5), lower stomach (1.5e4.0), or duodenum (7.0e8.5). Hydrogels capable of answering to these variations have special relevance due to their potential applications as responsive carriers able to deliver and release drugs or nutrients, among others [86e89], or as sensor and actuators [90,91]. pH-sensitive hydrogels are obtained from polyelectrolytes, polymers that present ionizable backbone or pendant groups, weak acidic or basic groups. The most common polymers used to develop this kind of materials are synthetic polymers such as

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Figure 12.5 Schematic illustration of the reversible self-folding of microgrippers in response to the temperature. They are composed on PPF segments in PNIPAAm-AAc layers. When temperature is above 36, the pNIPAAm-AAc layers exclude water and contracts which causes the grippers to open and close. When temperature is below 36, the PNIPAAm absorbs water and swells which causes the microgripper to open and to close in opposite direction. (Reproduced with permission from ACS J.C. Breger, C. Yoon, R. Xiao, H.R. Kwag, M.O. Wang, J.P. Fisher, et al., Self- folding thermo-magnetically responsive soft microgrippers, Applied Materials & Interfaces, 7(5) (2015) 3398e3405. (https://pubs.acs. org/doi/10.1021/am508621s). Copyright © (2015) American Chemical Society).

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poly(methacrylic acid, poly(acrylamide), poly(dimethylacrylamide), and their derivatives, among others, in addition to natural polymers including chitosan, hyaluronic acid, or alginate. The pH variation of the media triggers swelling process of the hydrogels by a change in the hydrodynamic volume or in the conformation of the polymer chains of the hydrogel due to the variation on the repulsion or attraction of the polar groups present in the matrix. This effect depends on the nature of the pH-responsive polymer, those hydrogels containing anionic polymer usually present in their structure weak acidic groups such as carboxylic or sulfonic acid groups. When the medium pH is below acid dissociation constant (pKa), the pendant groups are unionized and the hydrogel is in an unswollen state (Fig. 12.6(a)). When the pH is above the pKa, the acid groups are ionized leading to electrostatic repulsions between the negative charge of the pendant groups and the swelling of the network (Fig. 12.6(b)). However, some hydrogels under pH > 9e10 could unswell due to the charge screening effects of the counter ions (Naþ), which inherit the anion-anion repulsion [92e94]. The reverse phenomenon (a)

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