Self-Healing Properties of New Surface Treatments (EUROPEAN FEDERATION OF CORROSION SERIES)

European Federation of Corrosion Publications NUMBER 58

Self-healing properties of new surface treatments EFC 58

Edited by Lorenzo Fedrizzi, Wolfram Fürbeth & Fátima Montemor

Published for the European Federation of Corrosion by Maney Publishing on behalf of The Institute of Materials, Minerals & Mining

Published by Maney Publishing on behalf of the European Federation of Corrosion and The Institute of Materials, Minerals & Mining Maney Publishing is the trading name of W.S. Maney & Son Ltd. Maney Publishing, Suite 1C, Joseph’s Well, Hanover Walk, Leeds LS3 1AB, UK First published 2011 by Maney Publishing © 2011, European Federation of Corrosion The authors have asserted their moral rights. This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. Reasonable efforts have been made to publish reliable data and information, but the editors, authors and the publishers cannot assume responsibility for the validity of all materials. Neither the editors, authors nor the publishers, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused or alleged to be caused by this book. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming and recording, or by any information storage or retrieval system, without permission in writing from Maney Publishing. The consent of Maney Publishing does not extend to copying for general distribution, for promotion, for creating new works, or for resale. Specific permission must be obtained in writing from Maney Publishing for such copying. Trademark notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation, without intent to infringe. Maney Publishing ISBN-13: 978-1-906540-36-4 (book) Maney Publishing stock code: B810 ISSN 1354-5116 Typeset and printed by the Charlesworth Group, Wakefield, UK.

European Federation of Corrosion (EFC) publications: Series introduction

The European Federation of Corrosion (EFC), incorporated in Belgium, was founded in 1955 with the purpose of promoting European cooperation in the fields of research into corrosion and corrosion prevention. Membership of the EFC is based upon participation by corrosion societies and committees in technical Working Parties. Member societies appoint delegates to Working Parties, whose membership is expanded by personal corresponding membership. The activities of the Working Parties cover corrosion topics associated with inhibition, cathodic protection, education, reinforcement in concrete, microbial effects, hot gases and combustion products, environment-sensitive fracture, marine environments, refineries, surface science, physico-chemical methods of measurement, the nuclear industry, the automotive industry, the water industry, coatings, polymer materials, tribo-corrosion, archaeological objects, and the oil and gas industry. Working Parties and Task Forces on other topics are established as required. The Working Parties function in various ways, e.g. by preparing reports, organising symposia, conducting intensive courses and producing instructional material, including films. The activities of Working Parties are coordinated, through a Science and Technology Advisory Committee, by the Scientific Secretary. The administration of the EFC is handled by three Secretariats: DECHEMA e.V. in Germany, the Fédération Française pour les sciences de la Chimie (formely Société de Chimie Industrielle) in France, and The Institute of Materials, Minerals and Mining in the UK. These three Secretariats meet at the Board of Administrators of the EFC. There is an annual General Assembly at which delegates from all member societies meet to determine and approve EFC policy. News of EFC activities, forthcoming conferences, courses, etc., is published in a range of accredited corrosion and certain other journals throughout Europe. More detailed descriptions of activities are given in a Newsletter prepared by the Scientific Secretary. The output of the EFC takes various forms. Papers on particular topics, e.g. reviews or results of experimental work, may be published in scientific and technical journals in one or more countries in Europe. Conference proceedings are often published by the organisation responsible for the conference. In 1987 the, then, Institute of Metals was appointed as the official EFC publisher. Although the arrangement is non-exclusive and other routes for publication are still available, it is expected that the Working Parties of the EFC will use The Institute of Materials, Minerals and Mining for publication of reports, proceedings, etc., wherever possible. The name of The Institute of Metals was changed to The Institute of Materials (IoM) on 1 January 1992 and to The Institute of Materials, Minerals and Mining with effect from 26 June 2002. The series is now published by Maney Publishing on behalf of The Institute of Materials, Minerals and Mining. vii

viii

Series introduction

P. McIntyre EFC Series Editor The Institute of Materials, Minerals and Mining, London, UK EFC Secretariats are located at: Dr B. A. Rickinson European Federation of Corrosion, The Institute of Materials, Minerals and Mining, 1 Carlton House Terrace, London SW1Y 5DB, UK Mr M. Roche Fédération Européenne de la Corrosion, Fédération Française pour les sciences de la Chimie, 28 rue Saint-Dominique, F-75007 Paris, France Dr W. Meier Europäische Föderation Korrosion, DECHEMA e.V., Theodor-Heuss-Allee 25, D-60486 Frankfurt-am-Main, Germany

Contents

Series introduction

vii

Volumes in the EFC series

ix

1. 2. 3.

4.

5.

6.

Anticorrosive pigments in organic coatings Jörg A. Vogelsang

1

Self-healing anticorrosion coatings M. L. Zheludkevich, D. Raps, T. Hack and M. G. S. Ferreira

11

A review on the use of nanostructured and functional organosilane coatings modified with corrosion inhibitors as environmentally friendly pre-treatments for metallic substrates M. F. Montemor and M. G. S. Ferreira

39

Electrochemical study of cold rolled steel substrates pre-treated with silane films modified with CeO2 and TiO2 nanoparticles M. F. Montemor and M. G. S. Ferreira

65

Pyrrole-based silane primer for corrosion protection of commercial aluminium alloys M. Trueba and S. Trasatti

79

Sol–gel derived hybrid materials as functional coatings for metal surfaces J. Kron, K.-J. Deichmann and K. Rose

105

7.

Hybrid sol–gel/conducting polymer coatings: self-healing coatings for the corrosion protection of aerospace alloys R. Akid, M. Gobara and H. Wang 119

8.

Corrosion performance of nanoparticle-containing polyaniline films on AA3105 aluminium alloy O. Zubillaga, F. J. Cano, A. M. Cabral, P. J. Morais, I. S. Molchan, P. Skeldon, G. E. Thompson, T. Schmidt and M. Schem

134

Advances in the selection and use of rare-earth-based inhibitors for self-healing organic coatings S. J. Garcia, J. M. C. Mol, T. H. Muster, A. E. Hughes, J. Mardel, T. Miller, T. Markley, H. Terryn and J.H.W. de Wit

148

9.

10.

Corrosion inhibiting cerium compounds for chromium-free corrosion protective coatings on AA 2024 M. Schem, T. Schmidt, H. Caparrotti, M. Wittmar and M. Veith 184 v

vi 11.

12.

13.

14.

15.

16.

Index

Contents Hybrid Ce-containing silica methacrylate sol–gel coatings for corrosion protection of aluminium alloys M. Aparicio, N. C. Rosero-Navarro, S. A. Pellice, Y. Castro and A. Durán

202

Organosilicon plasma polymer coatings filled with Ce-based nanoparticles: characterisation of anti-corrosion properties D. Del Frari, J. Bour, J. Bardon, O. Buchheit, C. Arnoult and D. Ruch

220

Polypyrrole/aluminium flake hybrid pigments for corrosion inhibition of aluminium 2024 C. Vetter, X. Qi, A. C. Harper, S. V. Kasisomayajula and V. J. Gelling

238

Electrochemical behaviour of ZrO2 sol–gel films doped with corrosion inhibitors on AA2024 aluminium alloy F. Andreatta, L. Paussa, P. Aldighieri and L. Fedrizzi

262

Influence of the doping agent on the corrosion protection properties of polypyrrole grown on Al-2024–T3 A. Balaskas, I. L. Danilidis, I. Kartsonakis and G. Kordas

279

Self-healing coating with fluoro-organic compound for zinc A. Yabuki and R. Kaneda

293 306

1 Anticorrosive pigments in organic coatings Jörg A. Vogelsang Sika Technology AG, Tüffenwies 16, CH-8048 Zürich, Switzerland [email protected]

1.1

Introduction

It is widely known that corrosion causes direct and indirect costs amounting to about 6% of the gross domestic product (GDP) of a developed country. Organic coatings are the major protective measure accounting for up to 90% of expenditure on corrosion protection [1]. In order to increase the service life of organic coatings it has been good practice for many decades to add special pigments to the primer (the first coating layer on the substrate) which should increase the duration of corrosion protection. The purpose of these so-called active pigments is actively to influence the electrochemical processes taking place during the initiation of corrosion at the primer–substrate interface and at all later stages of degradation of the coating. It must be accepted that it is impossible to provide permanent protection to thermodynamically unstable low-alloyed steel. In all cases, it is only a question of time until the particular atmospheric conditions lead to irreversible water uptake, blister formation, delamination, crack formation and finally rust (Fig. 1.1 shows the ‘final stage’ of coating failure). Corrosion protection with organic coatings uses a wide range of different strategies and materials, e.g. multi-coat systems such as in Fig. 1.2 or high-build single coats of more than 1 mm thickness. Solvent-based materials still dominate for heavy duty corrosion protection such as outdoor applications in coastal zones but water dilutable coating materials or powder coatings are gaining increasing importance when controlled application conditions are available, e.g. in specialised paint workshops. In all branches of the industry, including the manufacture of cars and other forms of transportation (bus, truck, and rail), construction companies, shipyards and chemical plants, particular strategies for corrosion protection have been developed and all use, to some extent, organic coatings with active pigments. Some of them, for example, in steel construction, use coating systems similar to that shown in Fig. 1.2, where zinc particles act as an active pigment. Although active pigments have been used for decades, not much is known about the physico-chemical mechanism of their protective properties. This is true except for zinc dust primers, red lead oxide primers and chromate primers, where the basic mechanisms have been investigated and largely agreed by the scientific community. Unfortunately, red lead oxide and chromate pigments can no longer be used due to environmental concerns and safety at work restrictions (due mainly to their toxic and carcinogenic properties). The protection mechanisms for highly pigmented zinc dust primers have been quite well developed and will be presented below. 1

Anticorrosive pigments in organic coatings

1.1 Bicycle with steel frame which was transformed into its thermodynamically more stable state (iron oxides, rust) during weathering in the tidal zone of the Baltic Sea

1.2 Micrograph of a cross section of a typical coating system for heavy duty corrosion protection (e.g. C5M climate [2]), consisting of a zinc-rich primer, micaceous iron oxide (MIO) containing epoxy-based mid coat and MIO containing polyurethane-based top coat

2

3

Self-healing properties of new surface treatments

In recent years, it has become increasingly popular to talk about ‘self-healing’ in the context of active pigments and in fact, new concepts and materials are under investigation and development to complement the range of usable1 substances. These new concepts will briefly be explained in Section 1.8 of this paper. 1.2

Concepts for increased corrosion protection

Several reasonable concepts for increased corrosion protection are under discussion and can contribute to a longer lifetime for the coated product: • • • • • • • •

additional cathodic protection concealing of pores (multi-layer systems) blocking of pores and voids (by corrosion products) barrier formation selective reaction with oxygen improved wet adhesion (adhesion promoter) raising of pH value at the steel substrate precipitation of protective layer on the substrate.

In some cases, more than one of these concepts act together, simultaneously or in sequence. 1.3

Model description of protective properties of active pigments

As already mentioned, knowledge about how active pigments really work is very limited. Ideas and models are presented and discussed but with few hard facts for evidence. Passivation of the substrate by oxidising agents, inhibition by film formation of released molecules, precipitation of protective layers by interim partly soluble molecules or ions, enhanced barrier formation against corrosion stimulators and cathodic polarisation are the key models used to explain the protective properties of anticorrosive pigments. Passivation is obtained by oxidising agents which are able to stimulate the formation of a passive layer on the metal substrate. Chromate-containing pigments such as ZnCrO4 or SrCrO4 were used successfully, but nitrites can also have sufficient oxidising potential and sometimes red lead is also assigned to the group of passivators. In Fig. 1.3, the principal ideas of chromate reaction are summarised. Anodic oxidation of iron produces Fe2+ and the cathodic reduction of oxygen gives OH which react to form a passive layer of Fe(OH)2 on the steel substrate, yielding a well adherent oxide layer. In addition, chromates are reported to be attached to the iron oxides and incorporated into the passive layer. In some cases, mixed oxides of Fe(OH)2 and CrOOH are found which can explain the outstanding performance of chromate-based pigments.

1

‘usable’ in this context means user-friendly and environmentally acceptable by complying with environmental legislation, being reasonably priced and available in sufficient quantities.

Anticorrosive pigments in organic coatings

4

1.3 Sketch explaining the protective mechanism of chromate

1.4

Inhibition by film formation

Some organic molecules show excellent affinity for steel (iron oxide) surfaces and form mono- or multi-layers which are able to hinder the access of chlorides, oxygen and water to the substrate (Fig. 4). It must be emphasised that the adsorption of inhibitor molecules is reversible and that this adsorption stands in strong competition with corrosion stimulators. Therefore, this protective mechanism is likely to be the weakest. 1.5

Inhibition by precipitation

This mechanism is probably the most frequently occurring process; it applies to all phosphate- (including mixed phosphates), molybdate- and tungstate-containing

1.4 Organic inhibitor molecules have strong affinity for steel surfaces and compete for free adhesion sites with corrosion stimulators such as chlorides

5

Self-healing properties of new surface treatments

pigments and also to most metal oxide pigments such as red lead, zinc oxide and mixed oxides: •

• • • • •

phosphates • zinc phosphate • aluminium phosphate • chromium phosphate • multi-phase pigments borates • barium metaborate • zinc borophosphate boro-silicates molybdates, tungstates • zinc molybdate • calcium-zinc-molybdate ion exchanger (e.g. zeolite) metal oxide pigments.

In Fig. 1.5, the principal reactions are summarised. In all cases, partial solubility of the pigments is required. For example, phosphate ions are formed and iron ions are released from the steel substrate, which precipitate at the surface as insoluble iron phosphate. This precipitate protects the surface from further access of corrosion stimulators. The precipitate can also seal and block voids in the coating close to the surface. These active pigments also have good protective properties on imperfectly prepared surfaces with remaining traces of rust. Unfortunately, there is no general rule for the application of such partly soluble active pigments, because each generic type of binder (or worse still, each specific binder) requires comparative trials and tests to optimise the protective properties and to provide knowledge about the necessary amount of pigments in the formulation. There are not even any rules of thumb for this problem, just trial and error. 1.6

Barrier effect

This term describes the effect of mainly flake-like pigments in the coating formulation and is not limited to primers because even top coats can contain such flakes.

1.5 Film formation by precipitation of iron phosphate

Anticorrosive pigments in organic coatings

6

1.6 Barrier effect: diffusion hindered by flake-like pigments

During film formation of the freshly applied coating, the flakes are oriented mainly parallel to the substrate and develop a structure like that of a clapboard roof. Figure 1.6 illustrates the resulting effect: the diffusion of corrosion stimulators such as oxygen or water is hindered and slowed down through much longer diffusion paths in the coating. A contribution to the barrier effect can be seen when voids, pores and even small mechanical defects are sealed by, for example, corrosion products or precipitates of partly soluble active pigments (see section above). In particular, micaceous iron oxide (MIO) contributes to the outstanding protective properties of barrier pigments through a parasol effect; the organic binder beneath the flakes is shielded from harmful ultraviolet (UV) radiation and therefore the flakes remain on the surface although the binder above or around is degraded with time. Figure 1.7 illustrates the UV protection by MIO. Typical flake-like pigments are or can be formed by: • • • • • • 1.7

micaceous iron oxide (MIO) (see Figs. 1.8 and 1.9) glass mica aluminium zinc talcum. Cathodic polarisation

Galvanised steel has distinct advantages in corrosion resistance when compared to ordinary steel in atmospheric weathering conditions. The galvanic protection of the more noble steel by the less noble zinc is the reason for the better durability of zinc-coated steels. Therefore, the idea of incorporating large quantities of zinc particles into the primer formulation might lead to the illustrative picture of ‘cold galvanisation’ by zinc-rich primers. Actually, this picture is certainly incomplete and, if stretched too much, even misleading. In commercial products, the zinc content

7

Self-healing properties of new surface treatments

1.7 Glossy MIO flakes reflect ultraviolet (UV) light and protect the organic binder from degradation. This can be seen on older steel structures coated with MIO-containing paints, where a sparkling surface appears in the reflected light

varies and not all products reach the percolation level (overcritical pigmentation; the critical pigment volume concentration is given when not quite all of the particles are surrounded by binder and therefore have metallic contact) but also show very good corrosion protection even with a zinc content of about 60% by weight. The term ‘zinc-rich primer’ is used for formulations with about 90% plus of zinc dust particles [2]. Due to the metallic contact of the zinc particles in overcritical pigmented primers, the intrinsic conductivity is significantly higher than that of ordinary primers and this conductivity is required for the cathodic polarisation because the electrons from anodically dissolving zinc have to be transferred to the steel, where the cathodic part of the reaction takes place. As can be easily envisaged, the conductivity will vanish after a while because zinc corrosion products will form primarily in voids, but after further time, also between the particles, which leads to progressive isolation between the zinc particles. This isolation ends in a limited duration of cathodic polarisation, as can be seen in Figs. 1.9 and 1.10. 1.8 Example of ‘self-healing’ active pigments with multi-protective properties The effects of zinc particles in zinc-rich primers are better understood than many other pigments. As outlined above, to begin with, cathodic polarisation is the main protective effect, but after a while (perhaps some thousands of hours), the zinc particles have lost their conductive paths to the substrate and zinc corrosion products are formed in voids or pores. At this stage, zinc dust particles contribute to the barrier effect via cementation of zinc oxide or carbonate. It can be concluded that zinc dust offers at least two different mechanisms which act together in the protection of the substrate. With some justification, this mechanism can be called self-healing. Similar multi-protective properties are found when investigating red lead oxide. Although now forbidden in Europe for most applications, it offers outstanding corrosion protection: first, it acts as a passivator, second, it forms lead soaps with the

Anticorrosive pigments in organic coatings

8

1.8 MIO, top: optical micrograph of particles, thin plates appear red, thick plates are black. Bottom: micrograph of a polished cross section in bright field illumination, the MIO flakes reflect light and show parallel orientation

binder which precipitate at the substrate and third, it forms insoluble salts with the corrosion stimulators chloride and sulphate thus hindering depassivation by blocking access to the steel’s passive layer. New developments are aiming to take advantage of multi-protection pigments with up to four cooperative effects resulting in enhanced corrosion resistance of steel or aluminium structures when coated with primers containing such ‘smart’ self-healing pigments, e.g. see Fig. 1.11. Only one research group is cited here, but further activities are presented in other chapters of this book. Also, quite obviously, this book does not aim to summarise all activities worldwide. 1.9

Conclusions

This book presents some aspects of the current state of a new discussion about the mechanisms of self-healing and of new ways to work on such materials. This article

9

Self-healing properties of new surface treatments

1.9 Taken from Ref. 3: Change of potential with immersion time for steel coated with various zinc-rich epoxy primers

1.10 Taken from Ref. 4: Potential vs Time for a zinc-containing ethyl silicate coating with 18, 36, 46 and 63%Zn, respectively, under continuous immersion in 3.5% w/w sodium chloride solution

gives propaedeutic remarks to the field of active pigments and their use in corrosion protection by organic coatings. Unfortunately, most of the available information about anti-corrosive pigments, active pigments or self-healing pigments (all terms are synonymously used) cover only certain application conditions and practical examples for the use of the

Anticorrosive pigments in organic coatings

10

1.11 Illustration kindly obtained from Zheludkevich et al. [5]: triggering of inhibitor release from ‘smart’ self-healing nanocontainers

pigments. Knowledge about the protective mechanisms is not conclusive and, by far, is not exhaustive. Some mechanisms are more speculative than based on evidence. But these materials are used in large quantities with much success in two aspects: from a commercial point of view and – perhaps even more important – because of significantly improved corrosion protection. References 1. 2. 3. 4. 5.

NACE, Corrosion Costs and Preventive Strategies in the United States, www.nace.org and especially at http://events.nace.org/library/corrosion/Principles/CostPreventive.asp Paints & Varnishes – Corrosion Protection of Steel Structures by Protective Paint Systems. ISO 12944, parts 1, 2 and 5. ISO, Geneva, Switzerland, 1998. S. Feliu Jr, M. Morcillo, J. M. Bastidas and S. Feliu, J. Coat. Technol., 65, 826 (1993) 43–48. M. Walsh and J. D. Scantlebury, J. Corros. Sci. Eng. ISSN 1466-8858, Volume 2 Paper 24, submitted 15th October 1999. http://jcse.org/volume2/paper24/v2p24.php M. L. Zheludkevich, D. G. Shchukin, K. A. Yasakau, H. Mohwald and M. G. S. Ferreira, Chem. Mater., 19 (2007), 402–411.

2 Self-healing anticorrosion coatings M. L. Zheludkevich and M. G. S. Ferreira Department of Ceramics and Glass Engineering, CICECO, University of Aveiro, Aveiro, 3810-193, Portugal [email protected]

D. Raps and T. Hack EADS Innovation Works, 81663 Munich, Germany

2.1

Introduction

As already mentioned, the huge economic impact of corrosion of metallic structures is a very important worldwide issue. Engineered structures which suffer from corrosion attack include bridges, pipelines, storage tanks, automobiles, airplanes, ships and offshore installations. Corrosion has an impact on our daily life, very often causing not only economic consequences but also safety concerns, sometimes even resulting in life-threatening situations. The application of coatings is the most common and cost effective method of improving the corrosion resistance and consequently the durability of metallic structures. The main role of a coating in corrosion protection is to provide a dense barrier against corrosive species. However, defects appear in the protective films during operation of the coated structures, allowing direct access of corrosive agents to the metallic surface. The corrosion processes develop faster after disruption of the protective barrier. Therefore, active ‘self-healing’ of defects in coatings is necessary to provide long-term protection. The term ‘self-healing’ in materials science means self-recovery of the initial properties of the material following damage caused by the external environment or internal stresses. The same definition is also applicable to functional coatings. However, a partial recovery of the main functionality of the material can also be considered as a self-healing ability. Thus, in the case of corrosion protective coatings, the term ‘self-healing’ can be interpreted in different ways [1,2]. The classical understanding of self-healing is based on the complete recovery of the coating functionalities due to real healing of the defect, which restores the initial integrity of the coating. However, the main function of anticorrosion coatings is protection of a metallic substrate against environmentally-induced corrosion attack. Thus, it is not obligatory to recover all of the properties of the film in this case. The hindering of corrosion activity at a defect by the coating itself, by whatever means, is enough to constitute self-healing, because the corrosion protective system recovers its main function, namely protection against corrosion, after being damaged. Both concepts of self-healing will be taken into account in the present work. Historically, most of the effective active corrosion protection systems for metals have contained chromates. However, the outstanding corrosion protection provided 11

Self-healing anticorrosion coatings

12

by the leachability and high oxidation potential of chromates is accompanied by environmental hazards and toxic risks. Inhalation, skin contact, and ingestion may allow chromates to enter the human body. The hexavalent chromium species can be responsible for DNA damage and cancer [3]. The development of non-chromate environmentally friendly active corrosion protective systems is, therefore, an issue of prime scientific and technological importance for various industries due to the significant gap between industrial needs and currently existing corrosion protection technologies. Recent developments in the area of new environmentally-friendly self-healing anticorrosion coatings will be reviewed briefly in this chapter. The first part focuses on coatings, which can heal the defects to recover the coating integrity or self-seal them mechanically by corrosion products. Another part is devoted to the coatings which can provide active corrosion protection of metal in the defective areas by inhibition mechanisms including promising results obtained recently by the present authors. 2.2 2.2.1

Reflow-based and self-sealing coatings Coatings with self-healing ability based on the reflow effect

An important focus of current research efforts is the development of new bio-inspired self-repairing materials. Many ‘natural’ materials are themselves self-healing composites [4]. The repair strategies of living organisms attract materials designers looking for low-weight structures with enhanced service life. These bio-inspired approaches do not completely imitate the real biological processes involved because the latter are too complex. Instead, the designers of self-healing materials try to combine traditional engineering approaches with biological self-healing mechanisms. Several self-healing polymer composites have been reported recently. An outstanding example is an epoxy-based system able to heal cracks autonomically as described by White et al. [5]. The polymer bulk material contains a microencapsulated healing agent that is released upon crack initiation. Then, the healing agent is polymerised after contact with the embedded catalyst, bonding the crack faces and recovering the integrity of the material. The original idea of hollow spheres was subsequently extended to hollow reinforcement fibres, used in fibre reinforced plastic, embedded in a liquid resin. The repair process in this case is triggered after impact loading of the material [4,6–8]. The hollow glass fibres ranging in diameter from 30 to 100 μm and with hollowness up to 65% can be filled with uncured resin systems that bleed into a damage site upon fibre fracture, as shown in Fig. 2.1. After being cured, they provide a method of crack blocking and recovery of mechanical integrity. An even more advanced approach was recently suggested employing a self-healing system capable of autonomous repairing of repeated damage events [9]. The substrate composite delivers the healing agent to the cracks in a polymer coating via a threedimensional microvascular network embedded into the substrate. Crack damage in the coating is repaired repeatedly mimicking a body/skin system. However, this approach cannot be used in the case of corrosion protective coatings since a microvascular network cannot be created in a metallic substrate. Transfer of the self-healing approaches used for bulk materials to coatings is very complicated since the self-healing system should be embedded in a thin submillimetre polymer layer. The idea of a coating capable of reflow-healing of defects

13 Self-healing properties of new surface treatments

2.1 Crushed-healing fibres located under the impact site viewed under (a) normal and (b) UV illumination. Healing resin bridging cracked interface viewed under (c) normal and (d) UV illumination [7]

Self-healing anticorrosion coatings

14

was first patented a decade ago [10]. The invention describes protective coatings with in-situ self-repair ability after being damaged by stressful environments or careless handling. The repair is achieved through microencapsulated polymerisable agents incorporated into the coating matrix, as in the case of the bulk composites described above. If the coating is damaged, the ruptured microcapsules release the film-forming components in the immediate vicinity of the damage. The fluid flows over exposed areas of the metal surface and fills any defects or cracks in the coating, recovering the protective barrier. One of the first works in this area was focused on incorporation of micro-vesicles which can release the healing agent under destructive mechanical or chemical impact. The addition of the microencapsulated polymerisable agent to the fusion bonded epoxy coating was discussed as an advanced approach to the design of a more damage-resistant film than the traditional ones [11]. Kumar et al. attempted to introduce different types of capsules loaded with coating repair compounds and corrosion inhibitors into commercially available paints [12]. The efficacy of ‘self-healing’ corrosion protection coatings with urea-formaldehyde and gelatin microcapsules (50–150 μm diameter) containing several types of filmforming agents has been studied. The microcapsules stay intact for a long time in the dry coatings, as shown in Fig. 2.2a, and are ruptured only by damage that released the core constituents to the defect. The chosen microcapsules are stable for more than 2 weeks in the paint formulations. However, the results of the experiments suggest that they should be mixed with paint preferably at the time of application. Moreover, the best results were obtained when the microcapsules were sprinkled as a discrete layer on top of a thin layer of previously applied primer and, then, a second layer of primer was deposited followed by a topcoat layer (see Fig. 2.2b). The corrosion protection performance in this case was shown to be far superior compared to the mixing of the microcapsules in the primer before its application. Accelerated corrosion tests of these experimental coatings based on ASTM D 5894 indicate that the incorporation of self-healing microcapsules into commercially available primers can remarkably reduce underfilm corrosion on steel enclosures for outdoor equipment. Sauvant-Moynot et al. suggested using a self-healing coating together with a cathodic protection system. Specific film-formers sensitive to pH and electrical field were introduced into the coatings applied on metal structures used under cathodic protection [13]. Dried water soluble and self-curable epoxy electrodepositable additions as fillers (30 wt.%) were used as organic film formers. A significant reduction in the current needed for cathodic protection was revealed demonstrating the selfhealing ability of the coatings under study. The barrier properties were significantly increased in comparison to scratched reference samples. The idea of reflow-healing of protective coatings has already found commercial realisation. Nissan has recently announced the ‘Scratch Guard Coat’ painting system which contains a newly developed high-elastic resin providing reflow in artificial scratches [14]. The new coating system is effective for about 3 years and is five times more resistant to abrasion caused by car-washing compared with a conventional clear paint. 2.2.2

Self-sealing protective coatings

The examples of self-healing coatings presented above are based on polymerisation of a healing agent in the defects recovering the barrier properties of the protective coatings. However, the barrier properties of a damaged coating can also be restored

15

Self-healing properties of new surface treatments

2.2 (a) Micrograph of gelatin capsules in polyurethane paint after 2 h; (b) optical micrograph of cross section of coating with microcapsules [12]

by simple blocking of the defects with insoluble precipitates. These deposits in cracks can originate from the reaction of a corrosive medium or corrosion products with components of the coating. However, the term ‘self-sealing’ seems to be most appropriate in this case. Sugama and Gawlik developed a poly(phenylenesulphide) (PPS) self-sealing coating for carbon steel heat exchanger tubes, used in geothermal binary-cycle power plants operating at brine temperatures up to 200oC [15]. Hydraulic calcium aluminate (CA) fillers containing monocalcium aluminate (CaO.Al2O3) and calcium bialuminate (CaO.2Al2O3) reactants as major phases were dispersed in the coating matrix. The decalcification–hydration reactions of the CaO.Al2O3 and CaO.2Al2O3 fillers, surrounding the defect, lead to the fast growth of boehmite crystals in the cracks. The block-like boehmite crystals (~4 μm in size) fill the cracks after 24 h effectively sealing them, as demonstrated in Fig. 2.3. The sealing of the scratch causes an increase by about two orders of magnitude, of the pore resistance of the coating, suggesting that

Self-healing anticorrosion coatings

16

2.3 SEM images coupled with EDX spectra for cleaved PPS coatings with 5 wt.% CA fillers before (top) and after (bottom) exposure for 24 h to CO2-laden brine at 200oC [15]

the conductive pathway for aggressive species is thoroughly blocked. Extension of the exposure time to 20 days results in a stable value of pore resistance meaning that the sealing of the cracks by boehmite crystals plays an essential role in the recovery of the protective function of a PPS coating. Another self-sealing approach was used by Hikasa et al. for creation of a self-repair ceramic composite protective coating [16]. Sodium-clay (hectorite) and silica multilayers were deposited using a spin coating technique. The hectorite is a swelling clay which expands cubically due to the reaction with water. When water penetrates to a crack in the clay/silica composite film, it reacts with hectorite causing it to swell with consequent blocking of the defect. The encapsulation strategy was also suggested for the preparation of a polymeric self-sealing coating [13]. Epoxy-amine microcapsules containing MgSO4 and ranging in size from 10 μm to 240 μm diameter were prepared by interfacial polymerisation in inverse emulsion. A commercial liquid epoxy-amine paint doped with microcapsules and applied to a steel substrate was tested under cathodic protection conditions. Magnesium sulphate was chosen as a healing agent since it can form insoluble Mg(OH)2 precipitations at high pH which can arise in paint defects due to cathodic processes. The idea of this work was to achieve the sealing of the defects by the insoluble hydroxide formed by magnesium ions released from the capsules. However, the authors did not succeed with this and self-sealing was not achieved in this case.

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Self-healing properties of new surface treatments

2.3

Self-healing coatings based on active corrosion protection

The examples reviewed above demonstrate coatings which are able to heal or seal defects by mechanical blocking via polymerisation or precipitation mechanisms. The physical integrity of the coating is partially recovered due to these processes. However, the main function of the protective coatings is to prevent corrosion of the metallic substrate. The partial blocking of the pathways for corrosive species very often does not mean effective hindering of corrosion. Corrosion can occur beneath the coating where defects have been healed or sealed if the electrolyte penetrated to the substrate when the defect was induced. Therefore, other strategies have been proposed to achieve active suppression of corrosion processes in defective areas. Coating systems based on an active corrosion protection mechanism can also be considered as ‘self-healing’ coatings since the main function of the protection system, namely corrosion protection, is recovered during operation. A short overview of different active corrosion protection strategies and self-healing coatings based on them is presented below. Conducting polymers have been mainly investigated as barrier films for the protection of iron and copper. It was proven that the films can reduce the corrosion rate of copper in neutral chloride solutions [17]. In the case of polyaniline, a significant dependence of the protection efficiency on the redox state was found [18]. Michalik has shown that conducting polymer coatings have a certain passivating effect on small defects; however, in the presence of larger defects, the coating is no longer able to provide passivation. The consequence is rapid coating degradation followed by delamination [19]. The main disadvantages of conducting polymers are the limited chemical stability and the lack of UV-resistance. However, when conductive polymer is added in the form of particles or clusters to an insulating matrix, its degradation can be slowed, conferring a long-term passivation effect. Another approach for active corrosion protection and self-healing in corroded areas is based on the use of chemical inhibiting species which can be released from the coating system hindering the corrosion activity. The corrosion inhibitor can be added to the different parts of the coating system since the corrosion protection coating is usually a complex system constituted by several layers. The inhibiting compounds can be incorporated in the pre-treatment layer, the primer or the top coat using different strategies. The component containing the corrosion inhibitor serves as a reservoir from which inhibitors may leach out during its service-life. The following part of the paper provides an overview of different strategies for the incorporation of inhibitor into the corrosion protective systems to achieve self-healing and active corrosion protection. 2.3.1

Protective coatings with inhibitor-doped matrix

The easiest way to introduce a corrosion inhibitor is simply to mix it into the coating formulation. However, this procedure can raise many problems if some important factors are not taken into consideration. Firstly, the inhibiting species are effective only if their solubility in the corrosive environment is in the right range. Very low solubility leads to a lack of active agent at the defect site and, consequently, to a weak self-healing ability. If the solubility is too high, the substrate will be protected but for only a relatively short time since the inhibitor will be rapidly leached out from the coating. Another disadvantage, which can appear due to the high solubility, is the

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18

2.4 Osmotic blistering of organic coatings as a function of the solubility of corrosion inhibitor pigments [20]

osmotic pressure which forces water permeation leading to blistering and delamination of the protective coating. Figure 2.4 clearly shows how the solubility of the corrosion inhibitor influences the adhesion of the coating on the metal surface. A high solubility of the inhibiting species dispersed in the matrix causes rapid stripping of the polymer film from the substrate [20]. Another important problem can appear when the corrosion inhibitor interacts chemically with the components of the coating formulation weakening the barrier properties of the final coating. Degradation of the barrier properties resulting from the addition of an inhibitor is the main problem hampering the development of active corrosion protection systems. Moreover, interactions between the inhibitor and the components of the coating can lead to a complete deactivation of its inhibiting activity. Therefore, direct dissolution of an inhibitor in the formulations of organic coatings is not used in practice. The situation is different in the case of hybrid organosiloxane-based films. Several attempts to produce self-healing hybrid films doped with organic and inorganic corrosion inhibitors have recently been reported in the literature. Thin hybrid films are suggested as alternative pre-treatments or primers for different metallic substrates. The incorporation of inorganic or organic corrosion inhibitors into the hybrid films can significantly improve the corrosion protection properties. The results of several investigations have shown that the incorporation of cerium dopants into sol–gel coatings enhances the corrosion protection of aluminium alloys, magnesium alloys, galvanised steel, and stainless steel [21–30]. The corrosion protection properties of epoxy-zirconia sol–gel coatings with incorporated non-chromate inhibitors of Ce(NO3)3, NaVO3, and Na2MoO4 have been investigated by Voevodin et al. [31]. The results based on chemical analysis and electrochemical test methods showed that Ce(NO3)3 doped coatings exhibit the same grade of barrier as the undoped coatings [24]. The critical concentration of the cerium inhibitor is in the range of 0.2–0.6 wt.%. A higher concentration could lead to the formation of defects in the polymer network of the sol–gel film [23]. It was observed that sol–gel films with NaVO3 and Na2MoO4 did not display good barrier properties due to a decrease in the stability of the sol–gel network and coating delamination. However, improved corrosion protection was achieved by doping of molybdate anions within the siloxane

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Self-healing properties of new surface treatments

network of epoxy-silica hybrids, especially when the molybdate ions were added in a ‘bound’ form as an amine salt obtained from the reaction of the amine hardener and molybdic acid. The binding with the amine protects molybdate ions from interaction with the components of the hybrid matrix [32]. Organic inhibitors can also be incorporated into sol–gel matrices to improve corrosion protection of metallic substrates [33]. An additional inhibition effect was revealed when phenylphosphonic acid was introduced into a hybrid sol–gel film containing phenyl groups [34]. In several cases, the release of organic molecular species from the hybrid sol–gel matrix can be described by the pH-dependent triggered release mechanism [35]. The triggering of the desorption processes can provide an ‘intelligent’ release of the corrosion inhibitor only in places of local pH changes originating from localised corrosion processes. However, ionisable inhibitors show a significantly weaker effect than that of non-ionisable ones since the first are too strongly attached to the sol–gel matrix and, thus, cannot be released during corrosion [10]. Khramov et al. studied the corrosion protection properties of hybrid films with incorporated 2mercaptobenzothiazole and 2-mercaptobenzimidazole as corrosion inhibitors [36]. The SNAP (Self-assembled NAnophase Particle) sol–gel coatings are well suited to incorporate organic inhibitors. Inhibitor loaded SNAP films were evaluated by current density distribution maps for Al/Cu artificial defects on AA2024. The current peak after immersion for 3 h in NaCl solution was almost 25 times smaller in the case of the inhibited SNAP coating showing the inhibition of the corrosion process by the mercaptobenzimidazole molecules. Van Ooij and colleagues doped organosilane films with tolyltriazole and benzotriazole inhibitors. The organic inhibitor tolyltriazole, added to the silane film, improved the overall corrosion resistance of the AA2024-T3 alloy but did not impart a self-healing effect [21]. An organic corrosion inhibitor, tetrachloro-p-benzoquinone (chloranil), has also been incorporated into a hybrid organosiloxane/zirconia sol–gel matrix to improve corrosion protection [37]. The addition of a high content of chloranil leads to disorganisation of the sol–gel matrix and, consequently, to low corrosion protection. However, incorporation of lower concentrations of chloranil implicates homogeneous structures increasing the protective properties of the sol–gel coatings. The concept of ‘superprimers’ originates from the idea of combining pre-treatment and primer in one coating formulation. Such primers do not require a pre-treatment before paint application and can be applied directly to the bare metal. The addition of organofunctional silanes to conventional paint systems leads to good adhesion to both the substrate and to overcoats [38]. This study concerns the combination of an epoxy- and acrylate-based primer. The coating is formed from a water-based dispersion. As water evaporates after application, the silanes bond to the metallic surface whereas the acrylate and the epoxy form separated layers above it. Zinc phosphate incorporated in the hydrophilic acrylate-silane-zinc phosphate layer is able to leach out when the coating is scribed, creating a reservoir of zinc phosphate in the surrounding environment while the hydrophobic epoxy layer protects the acrylate and thereby the rest of the coated metal. A new approach for active corrosion protection of aluminium alloys using a metallic alloy film as a self-healing anticorrosive coating with multiple functionalities was recently suggested by Jakab and colleagues [39,40]. A new amorphous Al–Co–Ce alloy coating was synthesised and applied on the aluminium surface. At first, such metallic coatings can function by providing sacrificial anode-based cathodic protection. But the fact that the coating contains cerium and cobalt makes it also able

Self-healing anticorrosion coatings

20

to generate corrosion inhibiting species. The Co2+ and Ce3+ cations are released during anodic dissolution of the metallic coating. Dissolved cerium and cobalt cations can diffuse to the metallic substrate surface and form insoluble hydroxide precipitates covering the cathodic zones on the aluminium alloy. This work shows one of the first examples of a pH-controlled release of corrosion-inhibiting ions from an amorphous metallic coating where the pH change that results from the onset of corrosion triggers the inhibitor release. This inhibition strategy provides further corrosion protection beyond the traditional roles of barrier and sacrificial cathodic protection using a metal coating [39]. The level of protection conferred by these alloys was higher compared with conventional Alclad coatings in terms of both polarisation and maximum scratch size protected. The metallic amorphous coating can provide self-healing of scratches up to 2500 μm in width of exposed AA2024-T3 via a chemical inhibition mechanism [41]. An interesting example of a coating which can release inhibiting species on demand using a redox process as a trigger has been reported [42]. Galvanic reduction of 2,5-dimercapto-1,3,4-thiadiazole polymer in a conducting carbon paste releases its monomer anion. The monomer anion, in turn, exhibits very good inhibition efficiency for the cathodic oxygen reduction process. This chemistry forms the basis for a ‘smart’ self-healing material that releases an inhibitor when the material is coupled to a base metal, as in the case of a coating with a defect. No release would occur in the absence of the defect due to the lack of the reducing force of the base metal. Another inorganic Al2O3·Nb nanocomposite self-healing coating for corrosion protection of iron-based substrates was recently developed by Yasuda et al. [43]. The idea of using very thin, dense Al2O3 oxide layers on different metallic substrates as artificial passive films has been around since the beginning of the 1990s. However, films composed of only the oxide have no self-healing ability, resulting in rapid development of localised corrosion at the damaged site. Therefore, it is important for such films to possess self-healing abilities. Self-healing properties were achieved by introducing a metallic component into the oxide films. The oxidation of the metallic component can heal cracks in the composite films. The addition of metallic niobium to the oxide increases the self-healing ability but at the same time causes an undesirable increase in pinhole density. Therefore, composition-gradient films with the content of Nb increasing from the film surface to the substrate interface have been developed. For example, a composition-gradient Al2O3·Nb composite film with a niobium content ranging from 0 (top) to 96% (bottom) shows low pinhole defect density and high self-healing ability. In summary, the direct addition of a corrosion inhibitor to the coating formulation in some cases confers additional active corrosion protection and self-healing ability, especially when a metallic film is used as a sacrificial reservoir for corrosion inhibitors. However, in many cases, the inhibitor dissolved in the polymer coating causes weakening of the barrier properties and, consequently, of the overall corrosion protection. Therefore, other strategies of inhibitor introduction should be used to isolate the inhibitor from the coating components. 2.3.2 Self-healing anticorrosion coatings based on nano-/microcontainers of corrosion inhibitors A novel approach aimed at the development of a nanoporous reservoir for storing corrosion inhibitors at the metal/coating interface has been proposed by Lamaka

21

Self-healing properties of new surface treatments

et al. [44]. A porous titania layer on the surface of AA2024 aluminium alloy was developed using template-based synthesis, controllably hydrolysing titanium alkoxide in the presence of template agent. The reservoir is composed of titania nanoparticles which are self-assembled forming a cellular network that replicates the surface structure of the etched alloy. The formation of a network-like oxide-based structure with a highly developed surface area on the metallic substrate provides a great opportunity to load this layer with active substances. For this reason, the alloy with the applied porous nanostructured film was then immersed in an alcoholic solution of n-benzotriazole, which is a known corrosion inhibitor for AA2024 [45]. After loading the film with inhibitor, the substrate was coated with a hybrid sol–gel coating and then tested from the standpoint of corrosion protection. This novel pre-treatment resulted in enhanced corrosion protection in comparison with an undoped sol–gel film. The nanostructured titania reservoir layer covered with the hybrid film demonstrated well-defined self-healing ability, leading to effective long-term active corrosion protection [46]. The nanostructured porous character of the titania layer provided a very high effective surface area for the adsorption of the inhibitor. The adsorbed inhibitor was leached out from the porous pre-treatment layer to the defect, thus healing it. The developed surface, formed by the self-assembled layer, also offered good adhesion between the oxide and the sol–gel film due to the high contact area between the two phases. Moreover, the employment of this reservoir approach avoided the negative effects of the inhibitor on the stability of the sol–gel coating. New active multifunctional coatings should provide fast release of the active inhibiting species on demand within a short time of changes in the environment or the coating integrity. Recent developments in surface science and technology have yielded new concepts for the fabrication of self-healing coatings through the integration of nanoscale containers (carriers) loaded with active inhibiting compounds within existing conventional coatings. This approach leads to completely new coating systems based on ‘passive’ host–‘active’ guest structures. As a result, nanocontainers are uniformly distributed in the passive matrix keeping the active species in a ‘trapped’ state. This removes the possibility of excessive inhibitor leaching or pigment induced osmotic blistering that may occur if the inhibitor salts are too soluble or have too small a particle size. When the local environment changes or if a corrosion process starts at a coating defect, the nanocontainers respond to this signal and release the immobilised active material [47,48]. All concepts relating to inhibitor nanocontainers can be classified into two groups, namely: encapsulation with different types of shells; and immobilisation on the surface or inside carriers. In this section, the main approaches of inhibitor encapsulation and immobilisation on different carriers will be overviewed in terms of their applicability on self-healing anticorrosion coatings. A quite simple approach to inhibitor entrapment is based on the complexation of organic molecules by cyclodextrin [10,49]. This concept was originally suggested for controllable drug delivery systems [50]. Cyclodextrins are known complexation agents which can play the role of hosts forming inclusion complexes with various organic guest molecules that fit within the cyclodextrin cavities. Organic aromatic and heterocyclic compounds are usually the main candidates for the inclusion complexation reaction [51]. Several organic heterocyclic compounds are known inhibitors for various metallic substrates. Cyclodextrins can be effectively used for immobilisation of these species. Two organic corrosion inhibitors, 2-mercaptobenzothiazole (MBT) and 2-mercaptobenzimidazole (MBI), were selected by Khramov et al. [36] to be

Self-healing anticorrosion coatings

22

added to a hybrid sol–gel used for corrosion protection of AA2024 aluminium alloy. MBT and MBI were introduced in the sol–gel formulations as inclusion host/guest complexes with β-cyclodextrin. The hybrid films doped with corrosion inhibitors provided superior corrosion protection when compared to the undoped ones. Moreover, the coatings doped with inhibitors in the entrapped form outperformed those made by a simple addition. Figure 2.5 shows the appearance of an artificial scratch

2.5 Electrochemical impedance spectra for scribed hybrid coatings at different immersion times in dilute Harrison’s solution (A) without inhibitor and (B) with MBI/h-cyclodextrin complex. Inset: optical images of the samples after 4 weeks of immersion [10]

23

Self-healing properties of new surface treatments

and the impedance spectra on the sol–gel coated AA2024 after 4 weeks of immersion in dilute Harrison’s solution. The scratch remained almost shiny, without visible corrosion products, in the sample doped with the cyclodextrin-inhibitor complex. In contrast, white corrosion products covered the defect in the hybrid film directly doped with MBI. The higher values of impedance after such a long exposure to the corrosive electrolyte also confirmed the superior corrosion protection in the case of the film doped with immobilised MBI. Thus, formulations that contain b-cyclodextrin demonstrate superior corrosion protection properties because the complexation equilibrium results in slow release of the inhibitor and its continuing delivery to the corrosion sites followed by the self-healing of corrosion defects. However, complexation with cyclodextrin confers only the prolonged release of inhibitor without delivery on demand, such as in response to any external stimuli. Another entrapment concept is based on the use of oxide nanoparticles which can play the role of nanocarriers of corrosion inhibitors adsorbed on their surface. The oxide nanoparticles by themselves are known reinforcements for the coating formulations, as their addition leads to enhanced barrier properties [52–55]. The incorporation of nanoparticles into hybrid sol–gel formulations leads to an additional improvement in barrier properties, for example, due to the enhanced thickness and low crack sensitivity of such composites [55,56]. Moreover, additional active corrosion protection and a self-healing ability can be achieved when the oxide nanoparticles are doped with a corrosion inhibitor. Immobilisation of an inhibitor, e.g. Ce3+ ions, on the surface of the zirconia nanoparticles can be achieved during the synthesis of the sol by controlled hydrolysis of the precursors by a Ce-containing aqueous solution [57,58]. The resulting sol mixed with hybrid sol–gel formulation leads to nanocomposite hybrid coatings containing oxide nanocontainers of cerium ions. The high total surface area of the carriers, resulting from the small diameter of the nanoparticles (~4 nm as shown in Fig. 2.6), provides sufficient loading capacity

2.6 TEM image of a hybrid sol–gel film containing zirconia-based nanoreservoirs

Self-healing anticorrosion coatings

24

for the inhibitor. The prolonged release of inhibitor from the surface of oxide nanoparticles permits the enhanced long-term corrosion protection of AA2024 aluminium alloy compared to the case where the inhibitor is directly added to the sol–gel matrix. Moreover, the use of oxide nanocarriers prevents the negative effect of cerium ions on the hydrolytic stability of the hybrid sol–gel coating [58]. The inhibiting ions can also be immobilised at the surface of commercially available nanoparticles by simple adsorption in an inhibitor-containing solution. Silica- and ceria-based nanocarriers obtained by this method provide additional active corrosion protection to an organosilane coating applied to galvanised steel [59]. Another way of using the nanoparticle surface as an inhibitor carrier is to adsorb the inorganic corrosion inhibiting anions onto the particle surface by an ion-exchange mechanism [60]. The corrosion inhibitors are released from the particle surfaces by a subsequent ion-exchange with anions or cations (e.g. chlorides, sulphates, sodium ions) transported into the coating from the environment via water that penetrates through the coating. However, this mechanism can also lead to an undesirable release of inhibitor initiated by the presence of harmless ions in the surrounding environment, during, for example, cleaning, and not only by the ions causing the corrosion process. Organic corrosion inhibitors can also be immobilised at the surface of nanoparticles. Chemical anchoring of an organic inhibitor to aluminium oxyhydroxide nanoparticles through carboxylic bonds was employed in protective coatings for aluminium, copper, nickel, brass and bronze substrates [61]. Hydroxide ions generated from corrosion of these metals trigger the release of corrosion inhibitors from the particles at pH 9. Thus, the release of the inhibitor is initiated only by the corrosion process, preventing undesirable leakage of inhibitor from an intact coating during service. The high specific surface area of oxyhydroxide nanocarriers (at least 100 m2 gβ1) allows a higher quantity of corrosion inhibitor to be delivered into the damaged part of the coating. Even higher loading capacity can be achieved when porous fillers with hollow cellular structure are loaded with organic and/or inorganic inhibitors [62]. The hollow cellular structure material may be represented by diatomaceous earth, zeolite, or carbon. Zeolite particles are also attractive carriers because the cations in their structure are rather loosely held and can readily be exchanged for the inhibiting cations in the contact solution [63,64]. Inhibiting inorganic cations can also be incorporated as exchangeable ions associated with cation-exchange solids [65–67]. The advantage of this approach is that the cation exchange pigment is completely insoluble avoiding osmotic blistering [67]. Calcium(II) and cerium(III) cation-exchanged bentonite anticorrosion pigments are prepared by exhaustive exchange of naturally occurring Wyoming Bentonite. The cation exchange is carried out by repeated washing with aqueous solutions of cerium(III) chloride and calcium(II) chloride to produce bentonites containing 31 500 ppm exchangeable cerium(III) and 13 500 ppm exchangeable calcium(II), respectively. Bentonite clays are a form of montmorillonite and exhibit intrinsic cation exchange properties. They consist of stacked negatively charged aluminosilicate layers. The negative charge of these layers is compensated by the cations intercalated between the aluminosilicate sheets. The interaction between the sheets and the exchangeable cations is purely electrostatic. Therefore, the principal exchangeable cations (sodium(I) and calcium(II)) can be easily exchanged in the laboratory by passing a suspension of bentonite through a cation exchange column

25

Self-healing properties of new surface treatments

or by repeatedly washing the bentonite with a solution containing the desired cation. The calcium(II) and cerium(III) bentonite pigments have been dispersed (19% PVC) in polyester primer layers applied to pre-treated hot dip galvanised steel [67]. Salt spray test studies have shown that the calcium(II) and cerium(III) bentonite pigments are effective inhibitors of corrosion driven coating delamination at the cut edge of the polymer coated steel. Coatings doped with a Ce4+-modified ion-exchange pigment also demonstrate promising corrosion inhibition by cerium cations transported to the active defect sites on bare aluminium surfaces. The mechanism of corrosion inhibition using the Ce-doped cation-exchange pigment is presented in Fig. 2.7. The pigment keeps entrapped inhibiting cations

2.7 Schematic illustration of the potential mode of operation of a cationexchanged bentonite inhibitor pigment on a corroding galvanised surface [67]

Self-healing anticorrosion coatings

26

between the aluminosilicate sheets. The cations from the corrosive medium interact with the pigment particles, when the corrosive electrolyte penetrates to the defect in the coating, while exchanging the inhibiting cations. The Ce3+ ions diffuse, then, to the metallic surface and react with hydroxyl ions originating from cathodic oxygen reduction. As a result, the insoluble cerium hydroxide precipitates and blocks the cathodic sites on the metal surface, reducing the corrosion activity and healing the defect. However, as in the case of oxide nanoparticles and zeolites, the release of inhibitor here is based on the prolonged leaching and is not triggered by some specific corrosion related stimuli. A different situation arises when anion-exchange pigments are used to immobilise anionic inhibitors [68–71]. The release of inhibitor anions can be provoked in this case by aggressive corrosive chloride ions. The anion exchange pigment can play a double role absorbing the harmful chlorides and releasing the inhibiting ions in response. Even the ‘trap’ function alone can provide the additional anticorrosion effect in filiform corrosion tests when non-inhibited pigments with carbonates and sulphates are used in organic coatings [71]. The absorption of chlorides from an aggressive electrolyte in the vicinity of a defect reduces the aggressiveness of the corrosive medium, decreasing the rate of the corrosion processes. The use of an inhibiting anion-exchange pigment in addition can confirm the active feedback, conferring a self-healing ability. The layered double hydroxide compound hydrotalcite (HT) is a host–guest structure and can be used as an effective anion exchanger. The structure consists of a host, positively charged Al–Zn hydroxide layers, separated by layers of anions and water. The positive charge originates from the substitution of Al3+ on Zn2+ sites in the structure. In contact with an aggressive environment containing chlorides, an exchange reaction will occur. In this reaction, the inhibitor anions are released and the chlorides are adsorbed into the HT gallery. The exchange reaction is chemical in nature and is governed by the equilibrium constant for the reaction: HTInh + NaCl(aq) ȧ HTCl + NaInh(aq) [70]. In this expression, HTInh and HTCl refer to inhibitor and Clβ in the gallery of the HT compound, respectively. Buchheit et al. synthesised Al–Zn-decavanadate hydrotalcite pigments and added them to the epoxy-based coatings applied to AA2024 aluminium alloy. Results from corrosion tests illustrate an additional corrosion protection by the hydrotalcite pigment due to the decavanadate release accompanied by the uptake of the chloride ion in the exchange reaction [69]. The ability of HT to act as a generic anion delivery system creates the possibility of immobilising organic anionic inhibitors by intercalating them between hydroxide layers [72]. A hydrotalcite pigment doped with benzotriazolate, ethyl xanthate and oxalate anions was developed and tested in terms of inhibiting efficiency in respect of filiform corrosion on organically-coated AA2024. Inhibitor efficiency was found to increase in the order ethyl xanthate < oxalate < benzotriazole. However, it was not as efficient as chromate-based pigments and its practical usefulness remains unproven. Nevertheless, HT pigments doped with organic inhibitor anions appear worthy of further investigation as inhibitors for self-healing anticorrosion coatings. Various approaches to corrosion inhibitor immobilisation on different nano/ micro-particulated carriers have been discussed above. Another strategy is based on the use of different encapsulation techniques when a protective shell is created around a core containing inhibitor. The encapsulation of active healing agents for protective coatings has already been discussed in Section 2.2. A corrosion inhibitor can be encapsulated together with a polymerisable healing agent [12,13].

27

Self-healing properties of new surface treatments

Yang and van Ooij encapsulated triazole inhibitor using plasma polymerisation to produce PP-perfluorohexane and PP-pyrrole layers employing RF plasma discharge [73]. The plasma-treated triazole was used as a pigment in a water-based epoxy coating, slowly releasing the inhibitor and providing long-term corrosion protection. In both cases, the release of the inhibitor from the capsule is possible only when it is mechanically damaged. The damaged capsule releases the entire active agent very quickly in a non-controllable way. Microemulsion polymerisation is another inexpensive process well suited for the production of micron-sized capsules [74]. It is easily scalable and therefore easy to integrate in the manufacture of paint, providing the opportunity for fast commercialisation. The modification of current paint systems by microcapsules allows the rapid demonstration of emerging products and accelerates the replacement process of environmentally hazardous chromate pigments. Microemulsion oil-in-water polymerisation can provide a water-based suspension of microcapsules with a waterimmiscible organic core. Doping of the system with a corrosion inhibitor, which has low solubility in water, will lead to its preferential distribution in the organic phase. The formation of polymer shells around the organic microdrops forms microcapsules containing an organic core and loaded with corrosion inhibitor. A beneficial property of the liquid core material (diisopropylnaphthalene) used in this work is its ability to displace water. In the case of a defect in the organic paint film, the capsules are disrupted and release the water-displacing fluid, which removes the electrolyte and covers the scratch area. Furthermore, the inhibitors are delivered to active sites to passivate the surface and suppress the development and propagation of corrosion. Microcapsules loaded with mercaptobenzothiazole (MBT) as corrosion inhibitor were produced in this work and then introduced to commercial aeronautical primers. The undoped primer and primer doped with chromates were used here as references. Figure 2.8 shows SEM images of a conventional chromate-based coating system used in the aerospace industry and a new chromate-free coating system comprising a sol–gel coating as a pre-treatment and a capsule-loaded primer. Figure 2.8a shows the needle-like strontium chromate pigments in the conventional primer. Figure 2.8b shows the sol–gel coating as pre-treatment with a film thickness of about 6 μm and the primer with incorporated microcapsules. The capsules are uniformly distributed and adhere well to the primer matrix. No separation is observed on the primer/ microcapsule interface demonstrating the excellent compatibility of the primer and the capsules. Thermogravimetric analysis/differential scanning calorimetry (TGA/DSC) measurements (not shown) give evidence of the perfect enclosure of single capsules in the primer matrix. No physically trapped water (in between the capsules in the capsule paste) is released during the measurement because it evaporates with the diluent water of the water-based primer during drying [74]. The active corrosion protection potential of the capsule-loaded primer was compared to that of the conventional chromated one by a simple drop test. In this test, a scratch was produced on the coated surface and a drop of chloride-containing electrolyte was deposited on the scratch. Figure 2.9 shows exemplary images of the drop test samples with various paints on an anodic film as pre-treatment. This image reveals the unique properties of the chromate loaded primer. No pits were observed on the sample after 72 h exposure to the 3 wt.% NaCl electrolyte. The non-inhibited reference paint showed colouration after a short immersion time (24 h) and severe pitting corrosion. For the MBT-inhibited primers, higher loading degrees (MBT-10)

Self-healing anticorrosion coatings

28

2.8 SEM images after cryo-fracture of (a) a conventional coating system and (b) a microcapsule-containing coating system

of inhibitor led to better active protection in comparison to primer containing a lower amount of capsules (MBT-5). Scanning Vibrating Electrode Technique (SVET) measurements of the localised corrosion activity in artificial micro-defects supported the insights gained from the drop test. Figure 2.10 shows SVET maps of samples with a non-inhibited sol–gel coating as a pre-treatment coated with a non-inhibited primer, with primer doped with MBT enclosed in capsules, and with a chromate-containing primer. After 1

29 Self-healing properties of new surface treatments 2.9 Pictures of drop test samples (fully painted test specimen, anodic film as pre-treatment) on AA2024 unclad after 24 h, 48 h and 72 h exposure to 3 wt.% NaCl solution

Self-healing anticorrosion coatings

30

2.10 SVET maps of AA2024 unclad coated with a non-inhibited sol–gel coating and a non-inhibited primer (a), primer loaded with encapsulated MBT (b), and a chromate loaded primer (c) with an artificial scratch

month of exposure to 0.05 M NaCl, the non-inhibited reference sample showed a strong local current density in the scratch area, indicating cations produced in anodic places due to the oxidation of aluminium in the defect. The rather high current density indicated ongoing localised corrosion. The sample with encapsulated MBT showed only marginal cathodic activity in the scribe area. The minute currents detected emphasised the strong corrosion inhibition ability of the MBT-based primer. Almost no differentiation of cathodic and anodic areas was possible. The inhibition was even stronger than for the chromate-loaded primer where the anodic current density originating from pit formation and the cathodic current could be assigned to reduction of dissolved oxygen and upward flow of resulting OHβ ions [74]. A very interesting alternative, which allows controllable leaching triggered by corrosion related stimuli is the use of Layer-by-Layer (LbL) assembled shells. Nanocontainers with regulated storage/release of the inhibitor can be constructed with nanometre-scale precision employing the layer-by-layer deposition approach [75]. With such step-by-step deposition of oppositely charged substances (e.g.

31

Self-healing properties of new surface treatments

polyelectrolytes, nanoparticles, biomaterials) from their aqueous and non-aqueous solutions on the surface of a template material, LbL shells have been assembled and investigated as prospective materials for different applications [76,77]. LbL films containing one or several polyelectrolyte monolayers assembled on the surface of a sacrificial template possess the control of shell permeability towards ions and small organic molecules. The storage of corrosion inhibitors in the polyelectrolyte multilayers has two advantages: it isolates the inhibitor, avoiding its negative effect on the integrity of the coating, and provides an intelligent release of the corrosion inhibitor regulating the permeability of the polyelectrolyte assemblies by changing the local pH and humidity. The change of pH is the most preferable trigger to initiate the release of the inhibitor since, as is well-known, corrosion activity leads to local changes in pH near cathodic and anodic defects. Thus, a ‘smart’ coating containing polyelectrolyte containers can detect the beginning of the corrosion and start the self-healing process in the corrosion defect [47,48]. The possibility of creating such a smart self-healing anticorrosion coating based on LbL assembled nanocontainers was recently demonstrated by Zheludkevich and colleagues [78,79]. Silica nanoparticles were used as a template and benzotriazole as an organic corrosion inhibitor. The layer-by-layer deposition procedure was employed involving both large polyelectrolyte molecules and small benzotriazole ones. The initial SiO2 nanoparticles were negatively charged. Therefore, the deposition of the positive poly(ethylene imine) (PEI) was performed, in the first stage. Then, the deposition of the negative poly(styrene sulphonate) (PSS) layer was carried out. Deposition of the third inhibitor layer was accomplished in acidic media (pH 3) from a 10 mg mlβ1 solution of benzotriazole. The latter two deposition steps (PSS and benzotriazole) were repeated once to ensure a high inhibitor loading in the final LbL structure. Two PSS/benzotriazole bilayers have been identified to be the ideal number of deposited layers on the silica nanoparticles [79]. One bilayer is not sufficient for the self-healing effect of the final protective coating while three or more bilayers drastically increase aggregation of the nanocontainers during assembly and coating deposition. The benzotriazole content in nanocontainers is equal to 95 mg per 1 g of the initial SiO2 particles. The assembly process of such nanocontainers with incorporated corrosion inhibitors is schematically depicted in Fig. 2.11 [80]. Nanocontainers of

2.11 Schematic illustration of the procedure for benzotriazole loading of (A) SiO2 nanoparticulate containers and (B) halloysite nanotubes [80]

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this type cannot prevent spontaneous leakage; however, after 60 days of aging, the nanoparticulated reservoirs still contained benzotriazole in a quantity up to 80% of the initial inhibitor loading [80]. The sol of nanocontainers containing benzotriazole was mixed with a hybrid sol–gel formulation and then applied on the surface of AA2024 aluminium alloy [78,79]. The samples coated with the hybrid film doped with LbL nanocontainers demonstrated significantly enhanced performance in corrosion tests in comparison to an undoped sol–gel film or a film doped directly with free non-immobilised benzotriazole. After 2 weeks of immersion in a chloride solution, the nanoreservoir-containing film was still intact, while the sample coated with hybrid coating directly doped with benzotriazole showed extensive corrosion attack with many pits on the alloy surface (Fig. 2.12) [78]. To check whether the superior corrosion protection performance of a nanoreservoir-containing coating was related to its self-healing ability, SVET experiments were performed. A typical current map over an intact sol–gel film is depicted in Fig. 2.13a, illustrating the absence of local corrosion processes for both coatings. Artificial defects (200 μm in diameter) were formed on the surface of both coatings after 24 h of immersion in 0.05 M NaCl, as shown in Fig. 2.13b,f. Well-defined cathodic activity appeared at the induced defect on the alloy coated with the undoped hybrid film

2.12 AA2024 aluminium alloy coated with sol–gel film directly doped with (a) benzotriazole and (b) sol–gel film doped with LbL nanocontainers after 14 days of immersion in (a) 0.005 M NaCl and (b) 0.5 M NaCl (b) [78]

33

Self-healing properties of new surface treatments

2.13 SVET maps of the ionic currents measured above the surface of the AA2024 coated with undoped sol–gel pre-treatment (a, c, d, e) and with pre-treatments impregnated by LbL nanocontainers (g–i). The maps were obtained before defect formation (a) and for 4 h (c, g), 24 h (d, h) and 48 h (e, i) after defect formation. Scale units: I μA cmβ2. Scanned area: 2 mm × 2 mm [78]

(Fig. 2.13c). This activity became even more intense with longer immersion times (Fig. 2.13d,e). Significantly different behaviour was revealed after defect formation on the substrate coated with the hybrid film doped with nanocontainers. No corrosion activity appeared in this case 4 h after defect formation (Fig. 2.13g). Only after about 24 h did well-defined cathodic activity appear in the zone of the induced defect (Fig. 2.13h). The residual surface generated a cationic flow. However, the defect became inactive 2 h later and remained healed even after 48 h (Fig. 2.13i). One can see that local corrosion activity triggers the release of a portion of benzotriazole from the nanocontainers hindering the corrosion process in the defective

Self-healing anticorrosion coatings

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area. Such a ‘smart’ self-healing effect can originate from active feedback between the coating and the localised corrosion processes. The most probable mechanism is based on the local change of pH in the damaged area due to the corrosion processes. When the corrosion process is started, the pH value changes in the neighbouring area, which increases the permeability of the polyelectrolyte shells of the nanocontainers in a local area with the consequent release of benzotriazole. Then, the released inhibitor suppresses the corrosion activity and the pH value recovers, closing the polyelectrolyte shell of nanocontainers and terminating further release of the inhibitor, as shown schematically in Fig. 2.14. The promising results obtained on the LbL nanocontainers described above show the principal main opportunity of using this approach for intelligent self-healing coatings. However, nanocontainers with silica cores do not provide high inhibitor loading capacity. Porous cores are more promising. One of the prospective future containers may be industrially mined halloysite nanotubes. Halloysite is a viable and inexpensive nanoscale container ($4 per kg with a supply of 50 000 tons per year) for the encapsulation of biologically active molecules. The lumen of the halloysite can be used as an enzymatic nanoreactor [81]. The strong surface charge on the halloysite tubules has been exploited for designing nano-organised multilayers using the layer-by-layer method [82,83]. Halloysite is defined as a two-layered aluminosilicate, chemically similar to kaolin, which has a predominantly hollow tubular structure in the submicron range. The neighbouring alumina and silica layers create a packing disorder which causes them to curve. For most natural materials, the size of halloysite particles varies within 1–15 μm in length and 10–150 nm in inner diameter, depending on the deposits. The ſ-potential behaviour of halloysite particles is negative at pH

2.14 Scheme of the controllable release of the inhibitor from LbL nanocontainers and the ‘smart self-healing’ process [78]

35

Self-healing properties of new surface treatments

6–7 and similar to the surface potential of SiO2 with a small contribution from the positive Al2O3 inner surface (the chemical properties on the outermost surface of the halloysite nanotubes are similar to the properties of SiO2, while the properties of the inner cylinder core could be associated with Al2O3). However, at pH 8.5, the tubule lumen has a positive surface, promoting the loading of negative macromolecules and preventing their adsorption on the negatively charged outer surface. Halloysite nanotubes are capable of entrapping a range of active agents (within the inner lumen, as well as within void spaces in the multilayered aluminosilicate shell) followed by their retention and release. Both hydrophobic and hydrophilic agents can be embedded after an appropriate pre-treatment of the halloysite [84]. Inexpensive halloysite nanotubes as prospective nanocontainers for anticorrosion coatings with active corrosion protection have been demonstrated recently [80,85]. Halloysite nanotubes were loaded with the corrosion inhibitor 2mercaptobenzothiazole, which is partly soluble in water and well-soluble in ethanol or acetone, and then incorporated into a hybrid sol–gel coating. To prevent undesirable leakage of the loaded inhibitor from the halloysite interior, the outer surface of the 2-mercaptobenzothiazole-loaded halloysite nanotubes was modified by the deposition of several alternating polyelectrolyte multilayers (poly(allylamine hydrochloride) and poly(styrene sulphonate)) as shown in Fig. 2.11B. Another function of the polyelectrolyte shell is to provide the release of the encapsulated inhibitor in a way controlled by pH changes in the environment surrounding the halloysite nanotube [80], which will prevent the spontaneous leakage of the inhibitor and allow its release to be triggered by the pH changes directly in the corrosion pit. Halloysite nanocontainers showed very good upkeep characteristics – almost complete suppression of inhibitor release, with more than 90% of the initial loading retained inside the inner cavity. This can be explained by the geometrical restriction of the nanotubular container which is able to release the encapsulated material only through the polyelectrolyte-blocked edges with diameters of 20–50 nm. AA2024-T3 aluminium alloy was taken as a model metal substrate. The results of long-term corrosion tests demonstrated the superior corrosion protection performance of halloysite-doped hybrid sol–gel films compared to that of undoped coatings [85]. 2.4

Concluding remarks and outlook

An overview of different approaches for self-healing anticorrosion coatings has been presented here. Two completely different concepts of self-healing can be used for protective coating systems. The first is closer to the classical understanding of selfhealing and is based on mechanisms which allow recovery of the mechanical integrity of damaged coatings. However, a different approach, which is based on the active suppression of corrosion processes in defective areas, is also considered here as another self-healing concept. Looking to the future of self-healing anticorrosion coatings, the idea of a multilevel self-repair response seems to be the most promising. A multi-level self-healing concept would combine different damage repair mechanisms in the same coating system, which gradually activate in response to different environmental impacts and provide enhanced long-term protective properties. The different active components of the protective system must be able to respond to the different types and levels of impact imposed on the coating. A synergistic protective effect originating from the combination of different self-healing mechanisms can be achieved by incorporating

Self-healing anticorrosion coatings

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different types of nanocontainers in the same coating system. These nanocontainers can be added to the same polymer film or to different layers such as primer, clear coat or top coat depending on their functionality and the target of the active compounds. This approach will allow the creation of protective coatings which will adequately respond to environmental impacts, providing effective self-healing and a long service life. One can envisage a coating system which consists of a conversion coating with an active anticorrosion component, a primer doped with ‘smart’ nanocontainers of corrosion inhibitor and a top-coat with capsules of a polymerisable healing agent, providing outstanding corrosion protection properties and long-term performance without the necessity of repair. And finally, one of the most important issues, which limits the widespread use of self-healing coatings for different commercial applications, is the still relatively high cost of the suggested technologies. Only more extensive development of these approaches and further investment in the area will lead to lower prices opening the exciting possibility of seeing ‘smart’ self-healing coatings in our everyday lives. This review demonstrates the progress being made towards meeting these challenges and the feasibility of the self-healing idea for the development of new protective coating systems. References 1. W. Feng, S. H. Patel, M-Y. Young, J. L. Zunino and M. Xanthos, Adv. Polym. Technol., 26 (2007), 1–13. 2. W. Li and L. M. Calle, ‘Smart coating for corrosion sensing and protection’, in Proceedings of the US Army Corrosion Summit 2006 (Clearwater Beach, FL, 14–16 February 2006). 3. R. L. Twite and G. P. Bierwagen, Prog. Org. Coat., 33 (1998), 91–100. 4. R. S. Trask, H. R. Williams and I. P. Bond, Bioinspir. Biomim., 2 (2007), 1–9. 5. S. R. White, N. R. Sottos, P. H. Geubelle, J. S. Moore, M. R. Kessler, S. R. Sriram, E. N. Brown and S. Viswanathan, Nature, 409 (2001), 794–797. 6. G. Williams, R. Trask and I. Bond, Composites, 38 (2007), 1525–1532. 7. R. S. Trask, G. J. Williams and I. P. Bond, J. R. Soc. Interface, 4 (2007), 363–371. 8. J. W. C. Pang and I. P. Bond, Compos. Sci. Technol., 65 (2005), 1791–1799. 9. K. S. Toohey, N. R. Sottos, J. A. Lewis, J. S. Moore and S. R. White, Nat. Mater., 6 (2007), 581–585. 10. A. N. Khramov, N. N. Voevodin, V. N. Balbyshev and R. A. Mantz, Thin Solid Films, 483 (2005), 191–196. 11. D. G. Enos, J. A. Kehr and C. R. Guilbert, ‘A high-performance, damage-tolerant, fusion-bonded epoxy coating’, in Pipeline Protection Conference n°13 (1999) (Edinburgh, Scotland, 29 September–1 October 1999). 12. A. Kumar, L. D. Stephenson and J. N. Murray, Prog. Org. Coat., 55 (2006), 244–253. 13. V. Sauvant-Moynot, S. Gonzalez and J. Kittel, Prog. Org. Coat., (in press) 14. ‘Nissan develops world’s first clear paint that repairs scratches on car surfaces’, JCNN News Summaries, 5 December 2005. 15. T. Sugama and K. Gawlik, Mater. Lett., 57 (2003), 4282–4290. 16. A. Hikasa, T. Sekino, Y. Hayashi, R. Rajagopalan and K. Niihara, Mater. Res. Innov., 8 (2004), 84–88. 17. A. M. Fenelon and C. B. Breslin, Electrochim. Acta, 47 (2002), 4467. 18. V. Brusic, M. Angelopus and T. Graham, J. Electrochem. Soc., 144 (1997), 436. 19. A. Michalik and M. Rohwerder, Z. Phys. Chem., 219 (2005), 1547. 20. J. Sinko, Prog. Org. Coat., 42 (2001), 267–282.

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21. V. Palanivel, Y. Huang and W. J. van Ooij, Prog. Org. Coat., 53 (2005), 153–168. 22. L. S. Kasten J. T. Grant, N. Grebasch, N. Voevodin, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 11–15. 23. M. Garcia-Heras, A. Jimenez-Morales, B. Casal, J. C. Galvan, S. Radzki and M. A. Villegas, J. Alloys Compd., 380 (2004), 219–224. 24. N. N. Voevodin, N. T. Grebasch, W. S. Soto, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 24–28. 25. W. Trabelsi, P. Cecilio, M. G. S. Ferreira, K. Yasakau, M. L. Zheludkevich and M. F. Montemor, Prog. Org. Coat., 59 (2007), 214–223. 26. A. Pepe, M. Aparicio, A. Dur’an and S. Cer, J. Sol-Gel Sci. Technol., 39 (2006), 131–138. 27. M. F. Montemor and M. G. S. Ferreira, Electrochim. Acta, (2007), in press doi:10.1016/ j.electacta.2006.12.086. 28. A. M. Cabral, W. Trabelsi, R. Serra, M. F. Montemor, M. L. Zheludkevich and M. G. S. Ferreira, Corros. Sci., 48 (2006), 3740–3758. 29. M. F. Montemor, W. Trabelsi, M. Zheludevich and M. G. S. Ferreira, Prog. Org. Coat., 57 (2006), 67–77. 30. W. Trabelsi, E. Triki, L. Dhouibi, M. G. S. Ferreira, M. L. Zheludkevich and M. F. Montemor, Surf. Coat. Technol., 200 (2006), 4240–4250. 31. W. Trabelsi, P. Cecilio, M. G. S. Ferreira and M. F. Montemor, Prog. Org. Coat., 54 (2005), 276–284. 32. L. Mascia, L. Prezzi, G. D. Wilcox and M. Lavorgna, Prog. Org. Coat., 56 (2006), 13–22. 33. M. L. Zheludkevich, I. M. Salvado and M. G. S. Ferreira, J. Mater. Chem., 15 (2005), 5099–5111. 34. M. Sheffer, A. Groysman, D. Starosvetsky, N. Savchenko and D. Mandler, Corros. Sci., 46 (2004), 2975–2985. 35. A. J. Vreugdenhil and M. E. Woods, Prog. Org. Coat., 53 (2005), 119–125. 36. A. N. Khramov, N. N. Voevodin, V. N. Balbyshev and M. S. Donley, Thin Solid Films, 447 (2004), 549–557. 37. M. Quinet, B. Neveu, V. Moutarlier, P. Audebert and L. Ricq, Prog. Org. Coat., (2006), 1–8. 38. A. Seth, W. J. van Ooij, P. Puomi, Z. Yin, A. Ashirgade, S. Bafna and C. Shivane, Prog. Org. Coat., 58 (2007) 136. 39. M. A. Jakab and J. R. Scully, Nat. Mater., 4 (2005), 667–670. 40. M. A. Jakab, F. Presual-Moreno and J. R. Scully, Corrosion, 61 (2005), 246–263. 41. F. J. Presuel-Moreno, H. Wang, M. A. Jakab, R. G. Kelly and J. R. Scully, J. Electrochem. Soc., 153 (2006), B486–B498. 42. M. Kendig and P. Kinlen, J. Electrochem. Soc., 154 (2007), C195–C201. 43. M. Yasuda, N. Akao, N. Hara and K. Sugimoto, J. Electrochem. Soc., 150 (2003), B481–B487. 44. S. V. Lamaka, M. L. Zheludkevich, K. A. Yasakau, M. F. Montemor, P. Cecilio and M. G. S. Ferreira, Electrochem. Commun., 8 (2006), 421–428. 45. M. L. Zheludkevich, K. A. Yasakau, S. K. Poznyak and M. G. S. Ferreira, Corros. Sci., 47 (2005), 3368–3383. 46. S. V. Lamaka, M. L. Zheludkevich, K. A. Yasakau, R. Serra, S. K. Poznyak and M. G. S. Ferreira, Prog. Org. Coat., 58 (2007), 127–135. 47. D. G. Shchukin, M. Zheludkevich and H. Mohwald, J. Mater. Chem., 16 (2006), 4561–4566. 48. D. G. Shchukin and H. Mohwald, Small, 3 (2007), 926–943. 49. A. N. Khramov, N. N. Voevodin, V. N. Balbyshev and M. S. Donley, Thin Solid Films, 447–448 (2004), 549–557. 50. K. Uekama, F. Hirayama and T. Irie, Chem. Rev., 98 (1998), 2045–2076. 51. M. V. Rekharsky and Y. Inoue, Chem. Rev., 98 (1998), 1875–1917.

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52. Y. Chen, L. Jin and Y. Xie, J. Sol-Gel Sci. Technol., 13 (1998), 735–738. 53. J. Gallardo, A. Duran, I. Garcia, J. P. Celis, M. A. Arenas and A. Conde, J. Sol-Gel Sci. Technol., 27 (2003), 175–183. 54. A. Conde, A. Duran and J. J. de Damborenea, Prog. Org. Coat., 46 (2003), 288–296. 55. J. Malzbender and G. de With, Adv. Eng. Mater., 4 (2002), 296–300. 56. M. Zheludkevich, R. Serra, F. Montemor, I. Salvado and M. Ferreira, Surf. Coat. Technol., 200 (2006), 3084–3094. 57. M. L. Zheludkevich, R. Serra, M. F. Montemor, K. A. Yasakau, I. M. M. Salvado and M. G. S. Ferreira, Electrochim. Acta, 51 (2005), 208–217. 58. M. L. Zheludkevich, R. Serra, M. F. Montemor and M. G. S. Ferreira, Electrochem. Commun., 7 (2005), 836–840. 59. M. F. Montemor and M. G. S. Ferreira, Electrochim. Acta, 52 (2007), 6976–6987. 60. D. A. Pippard, ‘Corrosion inhibitors, method of producing them and protective coatings containing them’, US Patent 4405493, 20 September 1983. 61. R. L. Cook, ‘Releasable corrosion inhibitor compositions’, US Patent 6933046, 23 August 2005. 62. C. Schmidt, ‘Anti-corrosive coating including a filler with a hollow cellular structure’, US Patent 6383271, 7 May 2002. 63. S. Boston, J. Coat. Technol., 4 (2007), 167–175. 64. E. P. Eckler and L. M. Ferrara, ‘Anti-corrosive protective coatings’, US Patent 4738720, 19 April 1988. 65. R. G. Buchheit, S. B. Mamidipally, P. Schmutz and H. Guan, Corrosion, 58 (2002), 3–14. 66. H. N. McMurray, D. Williams, G. Williams and D. Worsley, Corros. Eng. Sci. Technol., 38 (2003), 112–118. 67. S. Bohm, H. N. McMurray, S. M. Powell and D. A. Worsley, Mater. Corros., 52 (2001), 896–903. 68. R. B. Leggat, W. Zhang, R. G. Buchheit and S. R. Taylor, Corrosion, 58 (2002), 322–328. 69. H. Wang, F. Presuel and R. G. Kelly, Electrochim. Acta, 49 (2004), 239–255. 70. R. G. Buchheit, H. Guan, S. Mahajanam and F. Wong, Prog. Org. Coat., 47 (2003), 174–182. 71. H. N. McMurray and G. Williams, Corrosion, 60 (2004), 219–228. 72. G. Williams and H. N. McMurray, Electrochem. Solid State Lett., 7 (2004), B13–B15. 73. H. Yang and W. J. van Ooij, Prog. Org. Coat., 50 (2004), 149–161. 74. D. Raps, PhD thesis, TU München, Germany, 2008. 75. G. Decher, J. D. Hong and J. Schmitt, Thin Solid Films, 210/211 (1992), 831–835. 76. S. L. Clark, E. S. Handy, M. F. Rubner and P. T. Hammond, Adv. Mater., 11 (1999), 1031–1035. 77. D. G. Shchukin, G. B. Sukhorukov and H. Mohwald, Chem. Mater., 15 (2003), 3947–3950. 78. M. Zheludkevich, D. G. Shchukin, K. A. Yasakau, H. Mohwald and M. G. S. Ferreira, Chem. Mater., 19 (2007), 402–411. 79. D. G. Shchukin, M. Zheludkevich, K. Yasakau, S. Lamaka, M. G. S. Ferreira and H. Mohwald, Adv. Mater., 18 (2006), 1672–1678. 80. D. G. Shchukin and H. Möhwald, Adv. Funct. Mater., 17 (2007), 1451–1458. 81. D. Shchukin, G. B. Sukhorukov, R. Price and Y. Lvov, Small, 1 (2005), 510–513. 82. Y. Lvov, R. Price, B. Gaber and I. Ichinose, Colloid Surf. Eng., 198 (2002), 375–382. 83. D. Kommireddy, S. Sriram, Y. Lvov and D. Mill, Biomaterials, 27 (2006), 4296–4303. 84. V. Luca and S. Thomson, J. Mater. Chem., 10 (2000), 2121–2126. 85. D. G. Shchukin, S. V. Lamaka, K. A. Yasakau, M. L. Zheludkevich, H. Möhwald and M. G. S. Ferreira, J. Phys. Chem. C, 112 (2007), 958–964.

3 A review on the use of nanostructured and functional organosilane coatings modified with corrosion inhibitors as environmentally friendly pre-treatments for metallic substrates M. F. Montemor Instituto Superior Técnico, ICEMS, DEQ, Av. Rovisco Pais 1049-001, Lisbon, Portugal [email protected]

M. G. S. Ferreira Instituto Superior Técnico, ICEMS, DEQ, Av. Rovisco Pais 1049-001, Lisbon; Department of Ceramic and Glass Engineering, University of Aveiro, 3810-193, Aveiro, Portugal

3.1

Introduction

Hybrid sol–gels obtained from organosilane solutions are very attractive methodologies for surface functionalisation of several metallic and non-metallic substrates. The formation of the silane layer is a very simple procedure, being achieved by dipping the metal in diluted alcohol or water-based solutions for a short period. The final result is a functional self-assembled coating that generally shows high stability and very good coupling properties. When applied over metallic substrates, the silane coating works like a ‘molecular bridge’ that promotes adhesion between the substrate and other organic layers, such as paints and adhesives. The silane/metal interface is characterised by the presence of chemically stable high strength bonds and the self-assembled coating confers good adhesion between the inorganic surface and organic polymers either in wet or dry environments and good barrier properties, which prevent moisture uptake and improved surface properties (scratch, wear, thermal and oxidation resistance). Because of these interesting properties, organosilicon-based coatings are used in many different fields. For example, silane coatings are applied within the micro array industry for coating of glasses and immobilisation of DNA molecules [1]. Silane coatings are used to modify the surface properties of Ti and stainless steel substrates, Co–Cr–Mo alloys used in biomedical applications [2–4], and to achieve controlled wettability of bioMEMS (biological microelectromechanical) systems [5]. Silanes can be used as fillers for the reinforcement of plastics, polymers and rubbers [6], for the development of sensors [5–7], for the modification of tribological and surface wear properties [7,8] and for application in electronics and optoelectronics [7,10–12]. The development of sol–gel technologies has been tremendous during recent years and the production of inorganic–organic hybrid materials has rapidly become a fascinating field. The chemistry and physical properties of the surface of a material can be radically altered and new functionalities can be introduced by using suitable organofunctional molecules. 39

The use of modified nanostructured and functional organosilane coatings

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The properties of thin organic coatings based on organosilane formulations have attracted the attention of corrosion researchers. Silanes were first used as adhesion promoters [13–22] but, due to their good barrier properties, they started to be applied as corrosion-resistant coatings. In 1986, Plueddemann et al. [14] reported that “high-performance polyfunctional silanes should find application in preparing corrosion-resistant composites for different applications”. Silanes were first applied on mild steel substrates to enhance their corrosion resistance in aggressive media [23–25]. Corrosion protection of aluminium alloys [26], steel, magnesium and copper was reported in the early 1990s [26–29]. Since then, many other works [30–53] have been published. There is a general agreement that silane-based coatings are adequate for the pre-treatment of a large number of metals and alloys, improving both adhesion and corrosion resistance. Silanes are, therefore, a suitable choice to replace anti-corrosion pre-treatments based on the use of hexavalent chromium. These pre-treatments were widely used and their success resulted from the very effective corrosion inhibition properties of Cr(VI) combined with low cost. However, chromates are very toxic [35] and the contact of this ion with skin or inhalation and ingestion causes DNA damage and cancer [36]. Furthermore, chromates cause an adverse environmental impact and consequently an important economic burden in what concerns environment protection. Recent regulations intend to eliminate the use of Cr(VI) containing compounds in anti-corrosion systems. Thus, research and development of environmentally friendly protective systems has been growing during recent years, and actually it constitutes one of the most challenging fields of modern surface engineering, presenting prime technological interest. Silanes are one of the most promising alternatives, being used already in many commercial formulations, especially in the coil coating industry. 3.2

Silane coatings

Silane molecules are typically trialkoxy esters (R1–Si–OR) where R is a methyl, ethyl or propyl group and R1 the main organic chain. The silicon atom existing in the silane molecule binds with three ester groups: –Si–(OR)3. If there is only one Si–(OR)3 end in the molecule, the silane is named monofunctional and if there are two Si–(OR)3 heads the silane is named bis-functional. Typically, a bis-silane molecule has the following structure: (RO)3–Si–(CH2)n–X–(CH2)n–Si–(OR)3 The main organic chain may contain additional functional groups (X) such as halides, sulphur, mercapto, amino, etc., which confer specific functionalities to the silane coating. The sulphur group, for example, confers higher hydrophobicity, whereas the amino group is more hydrophilic. These functional groups can be chosen to improve the organic compatibility that allows the organosilane to co-react with the functional groups existing in a paint or adhesive. This leads to the establishment of stronger bonds and therefore higher stability. 3.2.1

Formation of the silane coating

In the presence of water, the silane molecules are hydrolysed and the Si(OR)3 groups are converted into reactive silanol groups –Si–OH. This reaction releases an alcohol,

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the chemistry of which depends upon the alkoxy group present in the silane molecule. The formation of the self-assembled silane coating occurs according to a number of steps, among which the most important are [37]: 1) Hydrolysis:

R–Si(OR)3

R–Si(OH)3

2) Condensation

–Si–OH + HO–Si–

–Si–O–Si– + H2O

These are equilibrium reactions that proceed at rates that depend upon the functional groups existing in the silane molecule, water content, pH and temperature of the silane solution. All of these parameters and, additionally, the nature of the substrate and curing time have been investigated and found to play an important role in the formation of the protective surface coating on metallic substrates [38,39]. During dipping of the metallic substrate into the silane solution, the Si(OH)3 groups lead to the formation of bonds with the native metallic oxides/hydroxides present on the surface – the oxane bonds: M–O–Si. Ideally, this interfacial bridge must present very high bond strength and must prevent delamination of the coating from the substrate, even when water, oxygen and aggressive salts reach the interface. The silanol groups that could not bind with the substrate establish bonds among themselves, creating a network of silane molecules through Si–O–Si bridging. This process leads to the growth of the self-assembled silane coating that covers the entire substrate. The most important requirements of a silane coating used for corrosion protection of metallic substrates are: (i) high strength and stable chemical bonds between the silane and the substrate and between the silane and the paint or adhesive layer; (ii) uniformity and crack-free surface, providing an intermediate modulus that transfers stresses efficiently between the substrate and the coating; (iii) reversible hydrolytic bonds that allow for the reversible breaking of stressed bonds without loss of adhesion in the presence of aqueous media or chemical attack; (iv) thermal stability and high oxidation resistance; (v) effective barrier properties and (vi) adequate functionality. For optimum performance, it is very important that the chemistry of the organosilane and polymer be well matched and compatible. 3.2.2

Anti-corrosion properties

Organosilane-based coatings have been successfully tested as anti-corrosion pretreatments for aluminium alloys, steel, copper, magnesium and galvanised steel and the results generally demonstrate improved corrosion resistance [26–66]. The parameters that influence the reactions involved in the formation of the organosilicon-based coating play an important role in the anti-corrosion performance of the coatings. Franquet and colleagues [38,39], using spectroscopic ellipsometry, infrared spectroscopy and electrochemical techniques have shown that an increase in the concentration of the silane solution leads to thicker and more porous silane films. These authors also found that an additional curing step significantly decreased the film porosity, enhancing its anti-corrosion behaviour. Among the organosilane molecules that can be used as pre-treatments for corrosion protection of metallic substrates, bis-functional silanes have attracted special attention. Example of these silanes are the bis-[triethoxysilylpropyl] tetrasulphide (BTESPT) and bis-1,2-[triethoxysilyl] ethane (BTSE). Electrochemical measurements and accelerated corrosion tests, such as salt spray tests, show that these silanes provide enhanced corrosion protection of different metallic substrates [38–46]. This

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behaviour has been attributed to a barrier effect, which hinders the access of aggressive species towards the metallic interface. Some authors [40,48,49] also suggest that silane pre-treatment leads to the formation of a coating composed of two layers: an external organic-rich layer and an inner ‘interfacial layer’, which results from the formation of the oxane Si–O–M bonds (M denotes metal). Other silanes, such as aminosilanes, vinyl silanes and methacryl-silanes have been studied, showing good results for the protection of aluminium alloys and galvanised steel [50–53]. Concerning aminosilanes, there are authors reporting that the amino group can be hydrolysed leading to the formation of NH2+, which may attract chlorides and water to the film, accelerating its degradation [52]. Recent research work [52] suggested a mixture of bis-sulphur and bis-amino silanes for improved corrosion protection of aluminium alloys and galvanised steel substrates. This approach combines the hydrophilic nature of the amino silane with the hydrophobicity of the bis-sulphur silane that resulted in improved corrosion protection. Water-based mixtures of bis-amino silanes with vinyltriacetoxysilane also revealed good anti-corrosion performance [53]. The pre-treatment of the metallic substrate with conversion coatings before application of the silane coating formation is another way to improve the corrosion protection of galvanised steel. This procedure combines the anti-corrosion properties of the rare earth conversion coatings [44,45] with the properties of the silane films. However, in these studies, the anti-corrosion properties are, in most cases, a consequence of the good barrier effect created by the silane coating. Therefore, the corrosion performance of the pre-treated substrate will depend upon the silane layer thickness, uniformity, hydrophobicity and chemical stability. This makes the silane coating an ‘inert’ coating that cannot play any active role when the corrosion processes start to damage the surface. Furthermore, during corrosion attack, the cathodic processes release hydroxyl ions that increase the pH, causing decomposition of the silica network and consequently accelerated degradation and delamination of the silane coating. In order to overcome these limitations, a new challenge is imposed consisting of the modification of the bulk properties of the silane coating to make it more efficient with regard to the corrosion processes. This will increase the corrosion resistance of the metallic substrate and therefore the lifetime of the painted system. 3.3

More effective silane coatings through addition of active species

The modification of silane coatings with species able to inhibit corrosion introduces a new functionality to organic coatings: ‘active’ corrosion protection. This can overcome the typically inert nature of the silane film, making it more efficient in combating corrosion processes. The first approaches proposed in the literature [54,55] consisted of the addition of alumina or silica particles to improve the mechanical properties of the silane coating. The addition of these particles increased the impact, scratch and wear resistance. The corrosion resistance of aluminium alloys also seemed to increase with controlled amounts of particles [54]. This effect was attributed to the formation of silicate species that delayed corrosion activity. Silane films containing silica and formed under applied potential also revealed improved anti-corrosion behaviour when applied to aluminium substrates. In this case, critical silica contents were proposed [55]. The most recent trends consist of the addition of species with known anti-corrosion inhibition properties and the main goal is to achieve a self-healing ability. The term

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self-healing has been employed in many situations. The classical definition of selfhealing consists of the complete recovery of the functionality of a given system due to complete healing of the damaged parts, such as corrosion spots. However, nowadays, the term also includes the inhibition of corrosion activity at a defect in a coating, by a mechanism like the release of a corrosion inhibitor. If, afterwards, the coating or protective layer partially recovers its main functions, it is possible to state that self-healing was observed. A few authors have tried this approach. The modification of sol–gel coatings with organic inhibitors has demonstrated positive effects on their protective properties [56]. Environmentally compliant inhibitors such as Ce(NO)3, NaVO3 and Na2MoO4 have been incorporated into a Zr-epoxy sol–gel [57]. Some of these inhibitors, such as cerium nitrate did not damage the barrier properties, whereas others, such as sodium molybdate and sodium metavanadate reduced the barrier properties, promoting coating delamination. The corrosion behaviour of aluminium substrates treated with sol–gel systems containing cerium ions has demonstrated that cerium inhibits the corrosion processes [58]. It is also reported in the literature [59] that hybrid silica sol–gel coatings containing Ce3+ ions behave as conversion coatings on metallic zinc substrates. The anticorrosive performance of the Ce3+ ions entrapped in the hybrid silica sol–gel network occurs by an inhibitor effect and self-repairing mechanism (probably due to Ce(OH)3 precipitation). Recently, a new approach has been proposed in the literature for the formation of ‘smart’ self-healing anticorrosion coatings based on silica nanoparticles LbL-coated with polyelectrolyte molecules, which act as nanoreservoirs for corrosion inhibitors, incorporated in the hybrid sol–gel protective coatings [60]. The nanoreservoirs increase the long-term corrosion protection of the coated aluminium substrate and provide effective storage of the inhibitor and prolonged release ‘on demand’ to damaged areas, conferring active corrosion protection with a self-healing ability. 3.3.1

Addition of active ions

The addition of active ions, i.e. ions with well-known corrosion inhibition ability like cerium, zirconium or lanthanum to the silane formulations has been proposed previously by the current authors [61,62]. The aim of this procedure is to introduce a new functionality in the silane coating, making it more effective against corrosion. It is expected that these active species become trapped in the siloxane network, being released to the active corrosion sites, where they can develop its anti-corrosion ability. Furthermore, an improvement in the barrier properties of the coating is also expected due to its reduced porosity, increased thickness and decreased conductivity. Figures 3.1 and 3.2 show the electrochemical impedance spectroscopy (EIS) results obtained on bis-[triethoxysilylpropyl] tetrasulphide (BTESPT) silane coatings modified with cerium nitrate or zirconium nitrate and deposited on galvanised steel substrates that were immersed in aggressive NaCl solutions, as described elsewhere [61]. During the first 12 h of immersion, the EIS spectra obtained for the modified systems, especially for the Zr-doped ones, were characterised by the presence of a capacitive response over the whole frequency range. The phase angle was close to −90°, revealing that the silane coating behaved very much like a capacitor. Such a response revealed the presence of a defect-free and highly protective coating on the galvanised steel surface. With increasing immersion time, the behaviour of the EIS

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3.1 EIS Bode plots obtained for galvanised steel substrates pre-treated with bis-[triethoxysilylpropyl] tetrasulphide silane solutions doped with Ce(NO3)3 during immersion in (A) 0.005 M NaCl and (B) 0.5 M NaCl solutions. From Ref. 61

45 Self-healing properties of new surface treatments 3.2 EIS Bode plots obtained for galvanised steel substrates pre-treated with bis-[triethoxysilylpropy] tetrasulphide silane solutions doped with Zr(NO3)3 during immersion in (A) 0.005 M NaCl and (B) 0.5 M NaCl. From Ref. 61

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spectra changed and a resistive response developed at low frequencies. Such changes indicated that the electrolyte could reach the metallic substrate, developing conductive pathways. The resistive plateau, is, therefore attributed to the resistance of the silane pores. The low frequency impedance values, which are an indicator of the coating performance, depended upon the dopant and NaCl concentration, as shown in Figs. 3.1 and 3.2. For Ce-doped coatings immersed in the dilute NaCl solution (0.005 M), the low frequency impedance values were around 108 V cm2 and remained nearly constant with time (Fig. 3.1A). For the higher concentration of NaCl (0.5 M), the impedance values were above 107 V cm2 initially, but then decreased, attaining values of around 106 V cm2 after 1 week of immersion (Fig. 3.1B). For the most aggressive chloride solution, the increase in the phase angle in the low frequency range, after 2 days of immersion, coincided with the onset of corrosion. In fact, at the end of the tests, some signs of corrosion activity could be observed on the pre-treated surface in the form of small pits [61]. For substrates pre-treated with the Zr-doped silane and immersed in the more dilute solution, the first signs of corrosion were observed after 1 week of immersion. However, for substrates immersed in the more aggressive solution, corrosion activity was detected earlier and after 1 week of immersion, the corrosion signs were more evident than those observed on substrates pre-treated with the Ce-doped coating [61]. The impedance results allowed information to be obtained on the protective barrier properties of the modified silane coatings. Thus, in the early stages of immersion, the Zr-doped coatings provided very good barrier effects, but, as soon as these effects deteriorated, they could not delay the corrosion activity as effectively as the coatings modified with cerium. This can be related to the more effective corrosion inhibition ability of the cerium ions. Compared to an unmodified silane coating, the impedance increased by about three orders of magnitude, revealing that the addition of the active ions leads to much more protective coatings. Atomic force microscopy (AFM) scans obtained on the cerium- and zirconiummodified silane coatings revealed a very uniform and nanostructured surface, free of cracks and other defects [62]. Scanning electron microscopy (SEM) measurements revealed that the thicknesses of the modified silane coatings were about 2–3 times higher than those of an unmodified silane, which partially explains the improved barrier properties. Modification of two different silane coatings with lanthanum and cerium ions has also been reported [62]. It was shown that the corrosion behaviour depended upon the dopant and silane molecule. Compared to unmodified BTESPT coatings, the addition of cerium led to an increase of more than two orders of magnitude in the coating resistance, whereas the addition of La led to an increase of about one order of magnitude (Fig. 3.3). For bis-1,2 [triethoxysilyl] ethane silane (BTSE) films, these differences were not so marked. The resistance values were closer, being slightly higher for the Ce-containing films (Fig. 3.4). The beneficial effect of cerium-modified silane coatings has also been evaluated for aluminium alloys (AA2024-T3) [64]. Figure 3.5 shows the impedance plots obtained for samples treated with the BTESPT silane solution modified with cerium nitrate. Results obtained for the same alloy treated with a Cr(VI) conversion layer as well as results for the unmodified alloy, are also shown in Fig. 3.5 for comparison purposes. It can be seen that the Ce-modified silane solution displayed the highest impedance values, twice those for the Cr(VI)-treated sample.

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3.3 EIS Bode plots obtained for untreated galvanised steel and for galvanised steel pre-treated with: non-doped bis-[triethoxysilylpropyl] tetrasulphide (BTESPT) silane solution and doped bis-[triethoxysilylpropyl] tetrasulphide (BTESPT) silane solutions. Spectra were obtained after 1 day of immersion in 0.005 M NaCl. From Ref. 62

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3.4 EIS Bode plots obtained for untreated galvanised steel and for galvanised steel pre-treated with: non-doped bis-1,2 [triethoxysilyl] ethane silane (BTSE) solution and doped bis-1,2 [triethoxysilyl] ethane silane (BTSE) solutions. Spectra were obtained after 1 day of immersion in 0.005 M NaCl. From Ref. 62

In order to understand better the improved performance of silane coatings modified with cerium ions, electrochemical impedance experiments were conducted on galvanised steel samples treated with silane solutions modified with different concentrations of cerium ions as reported elsewhere [63]. The concentration of cerium

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3.5 Impedance results obtained for aluminium substrates (AA2024-T3) treated with cerium-modified silane solutions and treated with a Cr(VI)-based conversion coating. Spectra for the reference silane and spectra for the unmodified alloy are also included. Samples were immersed in 0.1 M NaCl. For the silane coating modified with cerium, spectra were obtained after 1 and 8 days of immersion. For the chromate treated samples, spectra were obtained after 1 and 7 days. For the reference silane and unmodified alloy, spectra were obtained after 1 day of immersion. From Ref. 64

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ions in the solution used to prepare the silane pre-treatment ranged from 1×10–4 M to 1×10–1 M. It was shown that the silane coating protective properties, i.e. coating capacitance and coating resistance, calculated after numerical simulation of the experimental impedance results, were strongly affected by the cerium concentration (Fig. 3.6). The most effective corrosion protection performance was observed for a concentration of 1×10–3 M. The results indicated that higher concentrations of cerium ions could induce the formation of defects in the silane coating, decreasing the barrier properties. Furthermore, it was shown that this concentration could effectively inhibit the corrosion activity at artificial defects formed in the coating. For that purpose, a coated galvanised steel substrate was immersed in the aggressive NaCl solution and a defect was created on the surface as described elsewhere [63]. After defect formation, current density maps were obtained using the scanning vibrating

3.6 Evolution of the coating resistance and coating capacitance (high frequency time constant parameters) during immersion in NaCl solutions. Values were obtained by numerical fitting using an equivalent circuit composed of two R/CPE components. From Ref. 63

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electrode technique (SVET) (Fig. 3.7). These maps depict the anodic (positive) and cathodic (negative) current densities. Anodic currents result from zinc oxidation, whereas cathodic currents result mainly from oxygen reduction. The maps showed the presence of anodic activity over the defect and cathodic activity around the defect as expected (Fig. 3.7A). The SVET maps also showed that the cathodic sites changed during immersion and that after 3 days of immersion, all corrosion activity vanished (Fig. 3.7B). However, for the unmodified coatings, anodic activity was permanently observed. The results show that in the presence of cerium ions, the defect was healed. Therefore, self-healing ability could be introduced in this silane coating [63]. The ability of silane coatings modified with cerium ions to inhibit corrosion activity was also tested, with success, for magnesium alloys [65]. AZ31 Mg alloys coated with a silane layer containing cerium ions revealed improved corrosion performance in electrochemical impedance tests. Furthermore, SVET measurements also showed

3.7 SVET results obtained on the substrates pre-treated with the 1×10–3 M cerium nitrate doped bis-[triethoxysilylpropyl] tetrasulfide (BTESPT) silane solutions: 8 hours after defect formation (A) and 2 days after defect formation (B). Results were obtained during immersion in 0.005 M NaCl. Scan size was 1×1 mm. From Ref. 64.

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a significant delay of corrosion activity at induced defects. Figure 3.8 shows the SVET current density maps and plots of current density obtained on AZ31 samples treated with the Ce-modified silane film and with non-modified silane after formation of a defect [65]. The current densities obtained on the non-modified silane were more than one order of magnitude above those obtained in the presence of cerium. Furthermore, it was shown that corrosion activity at the defects did not continue and that both anodic and cathodic current densities decreased with longer immersion times (Fig. 3.8). A microscopic and analytical investigation of the Ce-modified silane coatings formed on the AZ31 alloy revealed that the Ce-containing systems were thicker than the non-modified ones [66]. It was observed that the thickness decreased during immersion; however, the thickness of the Ce-containing coating was generally higher than that of the non-modified coating. Chemical information on the composition of the silane coatings formed on AZ31 substrates was obtained by X-ray photoelectron spectroscopy (XPS) analysis [66]. For this purpose, two Ce-modified silane coatings and two unmodified coatings were deposited on the substrate. The Ce-modified coatings contained two different concentrations of cerium ions (1×10–5 M and 1×10–2 M) and the same concentration of silane (5% v/v). One unmodified coating was obtained from a solution containing 5% (v/v) of silane and, another one, from a solution containing 0.5% (v/v) of silane. Details of the experimental procedure can be found elsewhere [66]. The XPS analysis showed that the chemical composition of the two unmodified silane coatings was identical before immersion in 0.005 M NaCl (Fig. 3.9). However, the chemical composition of the silane coatings modified with cerium ions was slightly different. The presence of cerium ions resulted in an increase in the silicon and oxygen contents and in a small decrease in the content of carbon and sulphur. After 48 h of immersion, these coatings still revealed a higher amount of silicon and no traces of magnesium, as depicted in Fig. 3.9. Cerium was only detected in the coating prepared with the highest concentration of cerium ions [66]. The analytical results indicate that the presence of cerium induced an increase in the silicon content either in the ‘fresh’ silane coatings or aged silane coatings, probably because cerium promoted the hydrolysis of the silane molecules, i.e. the formation of reactive silanol groups and therefore the assembling of the film and the final content of silicon. Cerium ions can, ultimately, complex with the reactive silanol groups, being incorporated in the siloxane network. This contributes to thicker surface films that are more compact and more organised, as demonstrated by SEM and AFM measurements, too [66]. The reported results have shown clearly that the addition of ions directly to the silane coating leads to a significant improvement in the barrier properties on different metallic substrates: galvanised steel, AA2024 alloy and AZ31 Mg alloy. The corrosion processes observed in different metallic substrates are also significantly delayed when the silane coating is modified with cerium ions and self-healing ability can be achieved in magnesium and zinc substrates. 3.3.2

Addition of nanoparticles

The addition of nanoparticles as fillers for silane or sol–gel coatings or the synthesis of particles inside the coatings are very attractive procedures, since nanoparticles can impart very interesting and specific properties as well as new functionalities to the modified coatings. Generally, small oxide particles provide improved resistance to

53 Self-healing properties of new surface treatments 3.8 SVET results. First row depicts the results obtained for the unmodified bis-[triethoxysilylpropyl] tetrasulphide (BTESPT) silane coating: (a) SVET map obtained 15 min after defect formation; Plot of the current densities over a 24 h immersion period and image of the surface at the end of the test (24 h). Second row depicts the results obtained for the silane coating modified with cerium ions: (b) SVET map obtained 15 min after defect formation; Plot of the current densities over a 72 h immersion period and image of the surface at the end of the test (72 h). Currents are μA/cm2 and the scan area is 2×2 mm. Adapted from Ref. 65

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3.9 Quantification of the elements present in the surface of the silane films with and without cerium prior to and after 48 h of immersion. Results were obtained from XPS analysis after peak fitting. BS is the bis-[triethoxysilylpropyl] tetrasulphide silane. Different coatings were tested: two unmodified coatings obtained from silane solutions with concentrations of 0.5% v/v of silane (BS 0.05%) and 5% v/v of silane (BS 5%) and two coatings prepared with 5% v/v of silane and different concentrations of cerium ions (BS + Ce 10–5 M and BS + Ce 10–2 M). From Ref. 66

oxidation, corrosion, erosion, scratch and wear. Nanoparticles can also impart UV resistance, colour changes and weldability. Many efforts have been made to enhance the corrosion resistance of metallic substrates by using ZrO2 [67,68], CeO2 [69,70], SiO2 [71,72], Al2O3 [73], and other single or mixed oxides. Concerning the use of nanoparticles, more elaborate procedures can be adopted. For example, the oxide nanoparticles can be doped with ions, such as cerium ions [74,75] to introduce corrosion inhibition ability, as these ions improve the corrosion resistance of silane coatings. The nanoparticles can also be covered with polyelectrolyte shells filled with corrosion inhibitor (LbL assembled nanoreservoirs) [60]. However, these new approaches raise some important questions: Which nanoparticles should be chosen? How stable are these nanoparticles? Are they also active? How can the agglomeration of nanoparticles be avoided? How does the concentration of nanoparticles influence the properties? In previous works [76,77], the present authors demonstrated that, concerning corrosion protection, CeO2 nanoparticles were more effective additives than SiO2 nanoparticles. Furthermore, it was demonstrated that CeO2 nanoparticles, by themselves, provided a significant delay of corrosion activity. However, they do not provide self-healing ability. It was demonstrated, by electrochemical impedance spectroscopy, that the content of nanoparticles in the silane coating plays an important role in the corrosion protection performance of silane coatings, as shown in Fig. 3.10. The low frequency impedance values depended upon the type and concentration of nanoparticles. For films modified with 100 and 250 ppm CeO2 the low frequency resistance values were identical (~10 MOhm.cm2), but there was a strong drop, of about two orders of magnitude, for films modified with 500 ppm CeO2 and the impedance values approached those of an unmodified silane film.

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3.10 EIS spectra obtained for galvanised steel coupons treated with bis-[triethoxysilylpropyl] tetrasulphide (BTESPT) solutions modified with different concentrations of SiO2 and CeO2 nanoparticles. Spectra obtained after 24 h of immersion in 0.005 M NaCl. From Ref. 66

The films modified with SiO2 showed an identical trend. However, the highest impedance values (~5 MOhm.cm2) were observed for films containing 250 ppm SiO2 and the lowest ones for films with 500 ppm SiO2. The EIS results indicate that the concentration of nanoparticles has an important impact on the barrier properties of

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the silane coatings. The lowest concentration resulted in better barrier properties, whereas, an increase in the concentration of nanoparticles had a negative effect on the barrier properties [77]. This trend may be related to the fact that nanoparticles are likely to form agglomerates that can be released during electrolyte uptake, creating large defects that promote faster uptake of the aggressive solution and, therefore, increased corrosion activity. Silane coatings modified with 250 ppm CeO2 or SiO2 nanoparticles and silane coatings modified with the same concentration of nanoparticles that were previously dispersed in a cerium nitrate solution were also studied [76]. The previous activation of the nanoparticles (SiO2 or CeO2) with cerium ions has two main objectives: to improve corrosion inhibition ability due to the presence of cerium ions and, simultaneously, to reduce the agglomeration of nanoparticles due to stabilisation of surface charging by the cerium ions as demonstrated elsewhere [75]. The activation of the nanoparticles with cerium ions led to a large increase in the total impedance of the system. For example, after 1 day of immersion, the total impedance of the CeO2 plus cerium coatings formed on galvanised steel (Fig. 3.11) was more than one order of magnitude higher compared to that of a coating modified with CeO2 only (Fig. 3.10). A small increase in impedance values was also observed for films filled with SiO2. Based on the EIS measurements performed for galvanised steel coupons treated with a silane modified with 250 ppm of nanoparticles, the effect of the nanoparticles in the barrier properties of silane coatings could be ranked as follows [76]: CeO2+Ce ions > SiO2+Ce ions ~ SiO2 ~ CeO2 > Unmodified silane coating. SVET measurements performed on scratched samples showed that the anodic and cathodic currents were very low for coatings containing CeO2 activated with cerium ions. For more than 3 days of immersion, a scratched surface revealed very low current densities, suggesting a strong inhibition activity. The SVET results also showed that the anodic activity at the scratch was stronger in the SiO2-containing systems and significantly delayed in the coatings containing CeO2 (Fig. 3.12). The SiO2-filled films, scratched and immersed in the aggressive solution were completely damaged after 24 h. However, in the presence of cerium ions, the degradation process and corrosion activity were delayed for both systems (SiO2 or CeO2). Potentiodynamic polarisation curves obtained for the scratched sample after 6 h of immersion in 0.005 M NaCl (Fig. 3.13) showed that the corrosion potential of the unmodified silane film was around –0.92 V, being identical to that of hot dip galvanised (HDG) steel. The most positive corrosion potential (–0.76 V) was that of the film filled with CeO2 activated with cerium ions. The kinetics of the anodic processes were strongly affected by the presence of the fillers. After 6 h of immersion the highest anodic currents were measured for the unmodified silane film (curve 5) and for the SiO2 filled film (curve 4) and the lowest ones for the CeO2 and cerium film (curve 1), in good agreement with the SVET data. The films containing CeO2 and cerium revealed a narrow passivation range at approximately 50–100 mV above the corrosion potential. All of the samples revealed an increase in the current density for more positive potential values. However, this increase was noticed at nobler potentials for the films containing CeO2 (curves 1 and 2). The trends observed in the anodic curves show that the presence of CeO2 nanoparticles induced a strong polarisation of the anodic processes, probably due to the

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3.11 EIS spectra obtained for galvanised steel coupons treated with bis-[triethoxysilylpropyl] tetrasulphide (BTESPT) silane solutions modified with 250 ppm of CeO2 or SiO2 nanoparticles activated with cerium ions. Spectra obtained after 24 h of immersion in 0.005 M NaCl. From Ref. 66

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3.12 (a) SVET maps obtained for the scratched SiO2 filled films after 24 h immersion. (b) SVET maps obtained for the scratched CeO2 filled films after 72 h of immersion. (c) SVET maps obtained for the scratched CeO2 plus cerium filled films after 72 h of immersion. Solution is NaCl 0.005 M. Left columns: maps of ionic currents; right columns: surface images. Scan size: 2×2 mm. Currents units: μA cm–2. Adapted from Ref. 65

formation of a more stable and more protective surface film. This behaviour agrees with the literature, where it is reported that CeO2 and Ce2O3 improve the anodic passivation behaviour of stainless steels [78,79], causing a strong shift of the potential towards the anodic direction. Nanoparticles imparted higher coating resistances, probably because they reduce conductivity/porosity of the silane layer. Due to changes in the viscosity of the silane solution, thicker coatings are also likely to form [77]. Concerning the effects on corrosion activity, the beneficial effect of nanoparticles can be described as follows: when a scratch is made on the surface, there is the development of anodic (exposed sites) and cathodic activity (around the scratch) wherever oxygen or another oxidant is available; the main cathodic reaction is oxygen reduction with the production of

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3.13 Anodic polarisation curves obtained for scratched galvanised steel substrates treated with bis-[triethoxysilylpropyl] tetrasulphide silane containing CeO2 and SiO2 nanoparticles, which were or were not previously activated with cerium ions. Curves were obtained after 6 h of immersion in 0.005 M NaCl solution. (1) CeO2 + Ce ions; (2) CeO2; (3) SiO2 + Ce ions; (4) SiO2; (5) Unmodified silane; 6-HDG. From Ref. 76

hydroxyl ions. Simultaneously, under the increased pH conditions, which may attain values of 11 [80] that develop at the cathodic sites, the silica component of the silane coating initiates a decomposition process, releasing nanoparticles that may precipitate together with the zinc corrosion products. However, SiO2 is not as effective as CeO2. This is because the SiO2 nanoparticles are also susceptible to alkaline decomposition with the formation of an expansive silicate-based gel, which promotes the degradation of particles. On the other hand, CeO2 is very stable under alkaline conditions. Thus, when the silane film decomposes, CeO2 nanoparticles are released from the silane coating and may precipitate together with the zinc corrosion products, reinforcing their protective role. The presence of cerium species in the scratched zone was detected by scanning Auger mapping and X-Ray photoelectron spectroscopy as shown in Figs. 3.14 and 3.15, respectively. Auger line scans were obtained over a line, which crossed a scratch made on the coating [77]. The spectra obtained over the silane coating revealed the presence of carbon, oxygen, silicon and sulphur, as expected. Over the scratch, the signal of zinc and oxygen increased, revealing the presence of zinc corrosion products. The cerium concentration was very

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3.14 Scanning Auger line profiles obtained for the film modified with SiO2 activated with cerium ions. Spectra obtained after 6 h of immersion in 0.005 M NaCl

low; however cerium could be detected together with the zinc corrosion products formed in the scratched zone after 6 h of immersion (Fig. 3.14). XPS analysis was also performed on the scratched substrates after 6 h of immersion. The XPS spectra obtained on the coatings filled with CeO2 nanoparticles activated with cerium ions revealed the presence of two forms of cerium: Ce(III) and Ce(IV). The latter is mainly due to the presence of the CeO2 nanoparticles, whereas the former confirms the existence of Ce3+, which may precipitate as Ce(OH)3. For the coating modified with SiO2 activated with cerium ions, the Ce3d ionisation displayed

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3.15 XPS spectra of the Ce3d region for the films modified with (a) CeO2 activated with cerium ions and (b) SiO2 activated with cerium ions. Spectra taken after 6 h of immersion in 0.005 M NaCl

a very low signal-to-background ratio; the oxidation states were not defined, however, cerium could be detected. The addition of cerium ions to the dispersions of nanoparticles improves the protective performance and seems to create an important synergy with the nanoparticles, reinforcing the protective behaviour of the silane films. Furthermore, nanoparticles impart good barrier properties and CeO2 nanoparticles contribute to a significant delay in the corrosion activity. Therefore, the addition of nanoparticles to silane formulations can be an effective way to increase the corrosion performance of the treated system and CeO2 nanoparticles play an active role against the corrosion processes. 3.4

Final remarks

Pre-treatments based on the use of bis-sulphur silane solutions are very effective for corrosion protection of different metallic substrates. Successful results have been obtained for galvanised steel, aluminium alloys and magnesium alloys. The silane coatings provide a protective barrier that delays the penetration of corrosive species to the metallic interface. The modification of the bis-sulphur silane solutions with corrosion inhibitors and/ or nanoparticles can be an effective way to improve both corrosion protection and durability of the silane coatings. The concentration of additives needs to be controlled and optimised, since the excess of additives may adversely affect the barrier properties.

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Cerium ions have proved to be a very effective additive for silane coatings and clearly improved the corrosion protection of galvanised steel and aluminium alloys. These ions enhance the protective barrier properties and inhibit corrosion activity at induced defects. The presence of cerium ions seems to develop a self-healing effect, hindering the corrosion processes at defects. The presence of small amounts of silica or ceria nanoparticles in the silane pretreatment also improves the corrosion resistance of galvanised steel substrates. Ceria nanoparticles are very effective and work as active anti-corrosion inhibitors, being much more effective than silica. The previous activation of the nanoparticles with cerium ions seems to create a synergistic effect that improves the corrosion resistance of silane pre-treated metallic substrates. The results obtained for modified silane coatings are very encouraging and many other alternatives are foreseen. The additions of nanoreservoirs, such as nanoparticles or halloysites, nanoclays and carbon nanotubes filled with corrosion inhibitors are examples of promising alternatives to develop self-healing silane coatings, providing higher corrosion resistance and longer lifetimes. References 1. S. D. Conzone and C. G. Pantano, Mater. Today, March (2004), 20. 2. J. P. Matinlinna, K. Laajaltho, T. Laiho, I. Kangasnieni, L. V. J. Lassila and P. K. Vallittu, Surf. Interf. Anal., 36 (2004), 246. 3. J. Liu, D. Yang, F. Shi and Y. Cai, Thin Solid Films, 429 (2003), 225. 4. P. J. Halling and P. Dunnill, Biotechnol. Bioeng., 21 (1979), 393. 5. K. C. Popat, S. Sharma, R. W. Johnson and T. D. Desai, Interf. Anal., 35 (2003), 205. 6. C. Kirkland, Plast. Technol., 27 (1981), 24. 7. G. Schottner, Chem. Mater., 13 (2001), 3422. 8. U. Bexell, M. Olsson, M. Johansson, J. Samuelson and P. E. Sundell, Surf. Coat. Technol., 166 (2003), 141. 9. R. Zhou, D. H. Lu, Y. H. Jiang and Q. N. Li, Wear, 259 (2005), 676. 10. T. Osaka, N. Takano and T. Yokoshima, Surf. Coat. Technol., 169 (2003), 1. 11. J. Xu and R. M. Almeida, Mater. Sci. Semicond. Process., 3 (2000), 339. 12. A. C. Marques, R. M. Almeida, A. Chiasera and A. Ferrari, J. Non-Crystal. Solids, 322 (2003), 272. 13. L. J. Matienzo, D. K. Shaffer, W. C. Moshier and G. D. Davis, Polym. Mater. Sci. Eng., 53 (1985), 592. 14. E. P. Plueddemann, P. G. Pape and H. M. Bank, Polym. Plast. Technol. Eng., 25 (1986), 223. 15. T. J. Lin, B. H. Chun, H. K. Yasuda, D. J. Yang and J. A. Antonelli, J. Adhes. Sci. Technol., 5 (1991), 893. 16. S. R. Holmes-Farley and L. C. Yanyo, J. Adhes. Sci. Technol., 5 (1991), 131. 17. A. Sabata, W. J. Van Ooij and R. J. Koch, J. Adhes. Sci. Technol., 7 (1993), 1153. 18. S. E. Hörnström, J. Karlsson, W. J. Van Ooij, N. Tang and H. Klang, J. Adhes. Sci. Technol.,10 (1996), 883. 19. T. F. Child and W. J. Van Ooij, Trans. Inst. Metal Finish., 77 (1999), 64. 20. W. J. Van Ooij, D. Q. Zhu, G Prasad, S. Jayaseelan, Y. Fu and N. Teredesai, Surf. Eng., 16 (2000), 386. 21. L. J. Matienzo, F. D. Egitto and P. E. Logan, J. Mater. Sci., 38 (2003), 4831. 22. D. Wang, Y. Ni, Q. Huo and D. E. Tallman, Thin Solid Films, 471 (2005), 177. 23. M. Lein, J. Coat. Technol., 54 (1982), 63. 24. M. Alagar, N. Aristotle, V. Krishnasamy and V. Mohan, Br. Corros. J., 21 (1986), 102. 25. H. Leidheiser Jr, M. De Costa and R. D. Granata, Corrosion 43 (1987), 382.

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26. P. J. Zanzucchi, and J. H. Thomas, J. Electrochem. Soc., 135 (1988), 1370. 27. M. A. Petrunin, A. P. Nazarov and Y. N. Mikhailovskii, Protection of Metals (English translation of Zaschita Metallov), 26 (1991), 749. 28. J. Jang and H. Ishida, J. Appl. Polym. Sci., 49 (1993), 1957. 29. J. Hansen, M. Kumagai and H. Ishida, Polymer, 35 (1994), 4780. 30. F. Deflorian, S. Rossi and L. Fedrizzi, Electrochim. Acta, 51 (2006), 6097. 31. A. Frignani, F. Zucchi, G. Trabanelli and V. Grassi, Corros. Sci., 48 (2006), 2258. 32. G. Pan and D. W. Schaefer, Thin Solid Films, 503 (2006), 259. 33. J. Flis and M. Kanoza, Electrochim. Acta, 51 (2006), 2338. 34. T. Van Schaftinghen, C. Le Pen, H. Terryn and F. Hörzenberger, Electrochim. Acta, 49 (2004), 2997. 35. J. H. Osborne, Prog. Org. Coat., 41 (2001), 28. 36. R. L. Twite and G. P. Bierwagen, Prog. Org. Coat., 33 (1998), 91. 37. G. Tesoro and Y. Wu, in Silanes and Other Coupling Agents, 215, ed. K. L. Mittal. VSP, The Netherlands, 1992. 38. A. Franquet, C. Le Pen, H. Terryn and J. Vereckeen, Electrochim. Acta, 48 (2003), 1245. 39. A. Franquet, H. Terryn, P. Bertrand and J. Vereckeen, Surf. Interf. Anal., 34 (2002), 25. 40. W. J. van Ooij, D. Zhu, M. Stacy, A. Seth, T. Mugada, J. Gandhi and P. Puomi, Tsinghua Sci. Technol., 10 (2005), 639. 41. V. Subramanian and W. J. van Ooij, Corrosion, 54 (1998), 204. 42. W. J. van Ooij and G. P. Sundararajan, J. Corros. Sci. Surf. Eng., 2 (2001) paper 14. 43. G. P. Sundararajan and W. J. van Ooij, Surf. Eng., 16 (2000), 315. 44. M. F. Montemor, M. G. S. Ferreira, R. G. Duarte and A. M. P. Simões, Electrochim. Acta, 49 (2004), 2927. 45. M. F. Montemor and M. G. S. Ferreira, Surf. Interf. Anal., 36 (2004), 773. 46. A. Cabral, R. G. Duarte, M. F. Montemor, M. L. Zheludkevich and M. G. S. Ferreira, Corros. Sci., 47 (2005), 869. 47. W. Trabelsi, L Dhouibi, E. Triki, M. G. S. Ferreira and M. F. Montemor, Surf. Coat. Technol., 192 (2005), 284. 48. D. Zhu and W. J. van Ooij, Corros. Sci., 45 (2003), 2177. 49. W. J. van Ooij and D. Zhu, Corrosion, 57 (2001), 413. 50. A. M. Beccaria and L. Chiaruttini, Corros. Sci., 41 (1999), 885. 51. A. Beccaria, G. Padeletti, G. Montesperlli and L. Chiaruttini, Surf. Coat. Technol., 111 (1999), 240. 52. D. Q. Zhu and W. J. van Ooij, Electrochim. Acta, 49 (2004), 1113. 53. D. Q. Zhu and W. J. van Ooij, Prog. Org. Coat., 49 (2004), 42. 54. V. Palanivel, D. Q. Zhu and W. J. van Ooij, Prog. Org. Coat., 47 (2003), 384. 55. L. Liu, J.-M. Hu, J.-Q. Zhang and C.-N. Cao, Electrochim. Acta, 52 (2006), 538. 56. N. N. Voevodin, V. N. Balbyshev, M. Khobaib and M. S. Donley, Prog. Org. Coat., 47 (2003), 416. 57. L. S. Kasten, J. T. Grant, N. Grebasch, N. Voevodin, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 11. 58. A. Pepe, M. Aparicio, S. Ceré and A. Durán, J. Non-Cryst. Solids, 348 (2004), 162. 59. M. Garcia-Heras, A. Jimenez-Morales, B. Casal, J. C. Galvan, S. Radzki and M. A. Villegas, J. Alloys Comp., 380 (2004), 219. 60. D. G. Shchukin, M. Zheludkevich, K. Yasakau, S. Lamaka, M. G. S. Ferreira and H. Möhwald, Adv. Mater., 18 (2006), 1672. 61. W. Trabelsi, E. Triki, L. Dhouibi, M. G. S. Ferreira, M. L. Zheludkevich and M. F. Montemor, Surf. Coat. Technol., 200 (2006), 4240. 62. M. F. Montemor, W. Trabelsi, M. Zheludevich, M. G. S. Ferreira and P. Cecílio; Prog. Org. Coat., 57 (2006), 67. 63. W. Trabelsi, P. Cecilio, M. G. S. Ferreira and M. F. Montemor, Prog. Org. Coat., 54 (2005), 276.

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64. A. M. Cabral, W. Trabelsi, R. Serra, M. F. Montemor, M. L. Zheludkevich and M. G. S. Ferreira, Corros. Sci., 48 (2006), 3740. 65. M. F. Montemor and M. G. S. Ferreira, Electrochim. Acta, 52 (2007), 7486. 66. M. F. Montemor and M. G. S. Ferreira, Prog. Org. Coat., 60 (2007), 228. 67. W. Liu, Y. Chen, C. Ye and P. Zhang, Ceram. Int., 28 (2002), 349. 68. A. Balamurugan, S. Kannan and S. Rajeswari, Mater. Lett., 57 (2003), 4202. 69. A. R. Phani, F. J. Gammel and T. Hack, Surf. Coat. Technol., 201 (2006), 3299. 70. W. Xiao, Q. Guo and E. G. Wang, Chem. Phys. Lett., 368 (2003), 527. 71. D. C. L. Vasconcelos, J. A. N. Carvalho, M. Mantel and W. L. Vasconcelos, J. Non-Cryst. Solids, 273 (2000), 135. 72. J. Gallardo, A. Durán and J. J. de Damborenea, Corros. Sci., 46 (2004), 795. 73. H. M. Hawthorne, A. Neville, T. Troczynski, X. Hu, M. Thammachart, Y. Xie, J. Fu and Q. Yang, Surf. Coat. Technol., 176 (2004), 243. 74. M. L. Zheludkevich, R. Serra, M. F. Montemor, I. M. Salvado and M. G. S. Ferreira, Surf. Coat. Technol., 2000 (2006), 3084. 75. M. L. Zheludkevich, R. Serra, M. F. Montemor, K. A. Yasakau, I. M. Miranda Salvado and M. G. S. Ferreira, Electrochim. Acta, 51 (2005), 208. 76. M. F. Montemor and M. G. S. Ferreira, Electrochim. Acta, 52 (2007), 6976. 77. M. F. Montemor, P. Cecílio and M. G. S. Ferreira, ‘Analytical characterization of silane films modified with cerium activated nanoparticles and its relation with the corrosion protection of galvanised steel substrates’, Prog. Org. Coat., in press. 78. E. Stoyonova, D. Nikolova, D. Stoychev, P. Stefanov and T. Marinova, Corros. Sci., 48 (2006), 4037. 79. D. Nikolova, E. Stoyonova, D. Stoychev, P. Stefanov and T. Marinova, Surf. Coat. Technol., 201 (2006), 1559. 80. M. F. Montemor, W. Trabelsi, S. V. Lamaka, K. A. Yasakau, M. L. Zheludkevich, A. C. Bastos and M. G. S. Ferreira, Electrochim. Acta, 53 (2008), 5913.

4 Electrochemical study of cold-rolled steel substrates pre-treated with silane films modified with CeO2 and TiO2 nanoparticles M. F. Montemor Instituto Superior Técnico, ICEMS, DEQ, Av. Rovisco Pais 1049-001, Lisbon, Portugal [email protected]

M. G. S. Ferreira Instituto Superior Técnico, ICEMS, DEQ, Av. Rovisco Pais 1049-001, Lisbon; Department of Ceramic and Glass Engineering, University of Aveiro, 3810-193, Aveiro, Portugal

4.1 Introduction Cold-rolled steel is a material with a very high economic and technological impact on socio-economic activities worldwide. However, steel substrates are very prone to corrosion attack. In order to minimise corrosion damage, steel is generally protected by a paint system. The protective behaviour of this system is critical for improved structure lifetime and reduced maintenance costs. The anti-corrosion performance of the paint system can be strongly enhanced when the steel surface is treated with an adhesion promoter. Some of the most successful adhesion promoters are based on the use of organofunctional silane molecules. These are currently used as adhesion promoters for coatings, mastic, joints and adhesives in various materials. Generally, a monosilane molecule has the following structure: X3–Si–(CH2)nY, where X3 corresponds to hydrolysable alkoxy groups surrounding a silicon atom. Y, which may or may not be present, is a functional group, such as amino, halogen or sulphur. When the silane molecule has two hydrolysable alkoxy groups, it is denominated a bis-silane such as X3–Si–(CH2)nY(CH2)n–Si–X3. Silanes have been widely studied and their properties as adhesion promoters are well-established [1,2]. Silanes also attracted the attention of industry and researchers because some of them may provide surface films with very good barrier properties. Therefore, silane coatings may delay corrosion activity on several metallic substrates including aluminium, steel, zinc and magnesium [2–7]. However, this effect is based essentially on the good barrier properties of the thin hybrid coating formed on the surface. These coatings are chemically stable, homogeneous and prevent electrolyte uptake. Therefore, silane films protect the substrate by acting as a physical barrier, but once the aggressive species reach the metallic substrate, the silane coating is no longer able to mitigate corrosion activity. More recently, the need for more effective surface treatments has led to the development of modified silane formulations. The most successful approach consists of the addition of species with good corrosion inhibition properties, such as cerium ions [8–10], organic inhibitors [11] or micro- and nanoparticles [12–15]. The best results 65

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have been obtained for aluminium substrates (especially AA2024), and galvanised steel. However, little has been published about the pre-treatment of cold-rolled steel substrates with modified silane solutions. The present work follows earlier studies [2,5,7–9,12–14] and aims at investigating the electrochemical behaviour of cold-rolled steel substrates pre-treated with bis-sulphur silane films modified with CeO2 nanoparticles and TiO2 nanoparticles. The treated substrates were studied during immersion in NaCl solutions using electrochemical impedance spectroscopy (EIS) and the scanning vibrating electrode technique (SVET). The influence of the role of nanoparticles on the corrosion activity of steel substrates is also discussed. 4.2 4.2.1

Experimental Materials and solutions

Nanoparticles of ceria (∅ ~10–20 nm) and titania (∅ <100 nm) were obtained from Sigma–Aldrich. They were ultrasonically dispersed in silane solutions to obtain a concentration of 250 ppm of nanoparticles. The bis-[triethoxysilylpropyl]tetrasulphide silane (Sigma/Aldrich product) solution was prepared by dissolving 5%(vol/vol) of the silane in a mixture of methanol (90% v/v) and water. The silane solution was stirred for 1 h and was kept for a few days before use. The substrate was cold-rolled steel (CRS), provided by a steel producer. The coupons were degreased with acetone, washed with distilled water, etched in a solution of 10% HNO3 for 2 min at approximately 50ºC, washed twice with distilled water, then dried prior to immersion in the silane solution containing the nanoparticles for 10 s. Finally, the treated coupons were cured in an oven at 120ºC for 40 min. 4.2.2

Electrochemical techniques

The EIS measurements were carried out using a Gamry FAS1 Femtostat with a PC4 Controller Board. The experiments were performed at room temperature, in a Faraday cage, at the open circuit potential, using a three-electrode electrochemical cell, consisting of the working electrode (~3.15 cm2 of exposed area), a saturated calomel electrode (SCE) as the reference electrode and platinum as the counter electrode. The measuring frequency ranged from 105 Hz down to 10–2 Hz. The r.m.s. voltage was 10 mV. The EIS experiments were performed during immersion of the pre-treated substrates in solutions of 0.05 M NaCl for 1 week. Spectra were analysed using Z-view software. The potentiodynamic polarisation curves were produced using RADIOMETERVOLTALAB PGZ 100 equipment. The scan rate was 1 mV/s, in the anodic direction, starting from the open circuit potential. An electrochemical cell was used consisting of three electrodes, as described for the EIS measurements. The electrolyte was 0.005 M NaCl. The SVET measurements were performed using Applicable Electronics equipment, controlled by the ASET program (Sciencewares). The vibrating electrode was made of platinum–iridium covered with polymer, leaving only an uncovered tip with a diameter of 40–50 μm. The distance from the tip to the surface was kept at 200 μm, as previously optimised. The scanned area had dimensions of 2 mm × 1.8 mm. To evaluate the corrosion inhibition performance of the modified silane film, the samples

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were immersed for 1 h in 0.005 M NaCl. A dilute solution was used to slow down the corrosion processes and to investigate, in more detail, the corrosion inhibition ability (the same solution was used for the polarisation measurements). After this period, a circular defect was introduced on the surface using a needle and further measurements were taken periodically. The diameter of the defect was approximately 0.20 mm. 4.3 4.3.1

Results EIS results

The silane coatings formed on the steel coupons were studied using electrochemical impedance spectroscopy (EIS) to assess their barrier properties (capacitance and resistance) and to assess information about the electrochemical behaviour of the steel/coating interface. Figures 4.1 to 4.3 depict the electrochemical impedance spectra obtained on the steel substrates pre-treated with the silane coatings modified with CeO2 and TiO2 nanoparticles after different immersion times in 0.05 M NaCl. Spectra for the unmodified (blank) silane coating are included. The low frequency impedance values are an indicator of the corrosion protection afforded by the silane coating. The spectra obtained after 6 h of immersion gave impedance values that depended on the composition of the silane coating. The highest impedance values, around 50 MOhm cm2, were obtained for the substrates treated with the silane coating containing CeO2 nanoparticles. The impedance values for the TiO2-containing coating were lower, around 1 MOhm cm2 and the lowest impedance values were measured for the unmodified silane coating. After 1 day of immersion, all of the impedance values decreased; the highest values were those measured on coatings containing CeO2. After 2 days of immersion, the substrates revealed some signs of corrosion attack and the impedance dropped to values more than one order of magnitude below those measured after the first 6 h (Fig. 4.1). The unmodified silane coating was severely damaged and red rust was observed. The phase angle revealed the presence of two time constants which were more or less dependent on the coating composition. The high frequency behaviour can be attributed to the response of the silane coating. After 6 h of immersion, the phase angle for the unmodified coating showed a broad maximum at a frequency of around 100 Hz. This maximum was shifted to higher frequencies (approximately 1000 Hz) in the modified coatings. This shift accounts for the more pronounced barrier effect and, therefore, more protective layer. The time constant on the low frequency side (between 1 and 10 Hz) can be attributed to the charge transfer controlled phenomena at the silane/steel interface. After 1 day, this time constant became better defined for all of the substrates. The high frequency time constant was still well-defined for the coatings modified with CeO2. After 2 days of immersion, significant changes were observed. The unmodified coating revealed a time constant at very low frequencies that can be attributed to diffusion-controlled processes across the layer of porous corrosion products formed on the interface. The phase angle shifted to values around 10 Hz for the TiO2 and CeO2 films, revealing a predominant effect of the corrosion processes on the EIS response. A quantitative evaluation was performed by numerical simulation of the EIS spectra using an equivalent circuit composed of two RC networks, as proposed elsewhere [15]. The parameters RHF and CPEHF relate to the

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4.1 EIS spectra obtained on CRS substrates treated with modified silane films and with unmodified (blank) silane after 6 h of immersion in 0.05 M NaCl

resistance and constant phase element of the silane film, respectively. The parameters CPELF and RLF relate to the low frequency processes and represent the double layer CPE and charge transfer processes, respectively, at the silane/metallic substrate interface. The barrier properties of the silane coating are reflected in the development of the high frequency parameters (Fig. 4.4). The high frequency resistance suffered a large

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4.2 EIS spectra obtained on CRS substrates treated with modified silane films and with unmodified (blank) silane after 1 day of immersion in 0.05 M NaCl

drop during the first day of immersion. At the very early immersion stages (first 6 h), the lowest resistances values were determined for the unmodified coating. The modified coatings showed resistances above 1 MOhm cm2, decaying with time and approaching the values measured for the unmodified film. During the first day of immersion, the highest values were observed for the coatings modified with CeO2. The coatings modified with TiO2 showed values that decreased below those of the

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4.3 EIS spectra obtained on CRS substrates treated with modified silane films and with unmodified (blank) silane after 2 days of immersion in 0.05 M NaCl

unmodified film. The capacitances of the doped films were around 50 nF/cm2, one order of magnitude below those of the unmodified film and showed a large rise after the first 6 h of immersion, revealing the development of conductive pathways in the film. The thicknesses of the films modified with CeO2 and TiO2 were estimated by scanning electron microscopy (not shown), being in the range of 2.5 to 3 μm, whereas those for the unmodified silane films were in the range 1–1.5 μm. These differences

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4.4 Variation with time of the fitting parameters for the high frequency time constant of the EIS spectra for the different samples during immersion in 0.05 M NaCl

also indicate the best barrier properties of the modified silane films and also help to explain the high impedance values observed for these systems. The time dependence of the high frequency parameters show that the silane coating was very prone to electrolyte uptake and that deterioration of the barrier properties

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occurred mainly after 1 day of immersion. This can be explained by the fact that steel substrates are very reactive in NaCl solutions. As soon as the electrolyte reaches the substrate, there is development of corrosion activity accompanied by alkalisation at cathodic sites. This promotes degradation of the siloxane network and, therefore, delamination of the silane coating. In fact, the low frequency resistance suffered a gradual drop for the modified systems, indicating an increase in the corrosion activity at the silane coating/metallic substrate interface (Fig. 4.5). During the first 12 h of immersion, the resistance values of the modified coatings at low frequency were about one order of magnitude higher than those of the unmodified ones. However, after 1 day, all of the values became similar, revealing important attack of the iron substrate. Simultaneously, the low frequency CPE increased by more than one order of magnitude. The development of the EIS plots over time revealed that the addition of CeO2 or TiO2 nanoparticles led to beneficial effects. These are reflected in both the barrier properties and the low frequency resistances, especially for films modified with CeO2. The beneficial effects of the TiO2 addition were not as marked as those of CeO2 and were only observed during the initial hours of immersion. 4.3.2

Polarisation results

Potentiodynamic polarisation curves were recorded after 30 min of immersion in 0.005 M NaCl at the open circuit potential (Fig. 4.6). Although potentiodynamic polarisation experiments are not recommended as a way to determine kinetic parameters from highly resistive coated systems, in this case, the technique was applied to help explain the observed polarisation effects. The anodic curves showed that the corrosion potential (Ecorr) was ennobled in the presence of the silane film. For the untreated steel, Ecorr was around –0.58 V and for coupons treated with the unmodified silane, it shifted to values of around –0.37 V. The corrosion potential became even more positive in the presence of nanoparticles, being approximately –0.32 V in the presence of TiO2 and –0.25 V in the presence of CeO2. Compared to the unmodified silane coating, the addition of CeO2 nanoparticles increased the potential by more than 0.1 V, suggesting that these particles led to significant polarisation of the anodic reactions. No significant effects could be detected in the cathodic branch. The differences between the silane treated steel and the untreated steel were greater than 0.2 V, revealing that the silane film alone caused polarisation of the anodic reactions, but the effect became more marked in the presence of nanoparticles. Plotting the anodic curves (not shown) as a function of (E-Ecorr), the lowest current density values were measured in the presence of CeO2 nanoparticles. For example, for potentials in the range 100–200 mV above the corrosion potential, current densities for the system with CeO2 were only one-half of those for the unmodified coatings. The changes observed in the anodic curves can be explained on the assumption that a more protective layer of corrosion products is formed when silane coatings modified with nanoparticles are deposited on the steel surface. 4.3.3

SVET results

The ability of the nanoparticles to inhibit or delay the corrosion processes was assessed on samples pre-treated with the modified silane coatings after forming an artificial defect. The treated samples were immersed in 0.005 M NaCl for 1 h after

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4.5 Variation with time of the fitting parameters for the low frequency time constant of the EIS spectra for the different samples during immersion in 0.05 M NaCl

which a defect was created on the surface as described in the experimental section. The defects were formed using a needle with a diameter of approximately 0.2 mm. The reproducibility of these defects is an important point. Therefore, they were all made using the same procedure, although small differences in shape could occur due

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4.6 Anodic polarisation curves for the different samples during immersion in 0.05 M NaCl

to film rupture around the defect. The error related to reproducibility is over 10%, which is acceptable for differences in current of about one order of magnitude, as observed in this work. Figure 4.7 shows the maps obtained for the different films 30 min after formation of the defect. The unmodified film revealed anodic activity over the defect previously formed on the surface. The maximum anodic current density reached values of approximately 10 μA/cm2. Intense cathodic activity was observed close to the anodic area, where oxygen could reach the interface. For films containing nanoparticles, the currents were much lower, being below 1 μA/cm2 for both films. With time, there was a significant increase in anodic activity and an enlargement of the anodic area (Fig. 4.8) in the unmodified system and in the system modified with TiO2. After 1 day of immersion, these samples revealed strong corrosion attack and a large amount of red rust around the initial defect. However, this was not observed for the coating modified with the CeO2 nanoparticles. After 48 h, this system revealed current densities close to the initial values, suggesting that the corrosion process was significantly delayed, compared to the other systems under test. 4.4

Discussion

The incorporation of inorganic nanoparticles into an organic or hybrid matrix is envisaged as an approach of interest in many applications. The addition of nanoparticles can improve specific properties such as abrasion and wear resistance, barrier effects, UV shielding, high-temperature oxidation, electrical conductivity and corrosion resistance among others. Thus, the addition of nanoparticles can improve existing properties or can introduce new functionalities in organic or hybrid coatings.

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4.7 SVET maps obtained during immersion in 0.005 M NaCl 30 minutes after formation of an artificial defect. Currents units: μA.cm–2; scan size 0.2 × 0.2 mm. (a) blank silane film; (b) silane film modified with CeO2; (c) silane film modified with TiO2.

The interface between the nanoparticles and the organic matrix plays an important role in the final performance of modified coating systems. It is reported in the literature that the nanoparticles, when embedded in a silane matrix can bind covalently to the polymer matrix via the silanol groups. This has been demonstrated for several nanopowders functionalised with silanes [16–19]. CeO2 is a stable oxide that exists in the cubic fluorite structure (CaF2). However, it may form oxygen vacancies as well as interstitial defects, becoming reactive through charge-compensating defects in the oxygen sub-lattice. This makes these nanoparticles suitable, for example, for the adsorption of charged species such as metallic cations [20]. In neutral media, the surface of CeO2 may develop OH– or OH2+ groups [21,22]. Therefore, the hydrolysed nanoparticles can bind to the silane molecules via the silanol groups as proposed in the literature [21] and shown in the scheme presented in Fig. 4.9. The presence of surface defects in the TiO2 structure has also been reported [23] and reaction via the silanol groups is also likely to occur [23,24] as proposed in Fig. 4.9 for CeO2. The formation of such bonds enhances the adhesion of the nanoparticles to the siloxane network and may help to decrease porosity and conductivity due to filling of defects and voids in the reticulated siloxane network. This effect may lead to the improved barrier properties observed in the EIS measurements during the early stages of immersion.

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4.8 SVET maps obtained during immersion in 0.005 M NaCl after formation of an artificial defect. Left side: current maps; Right side: image of the surface. Current units: μA.cm–2; scan size 2 × 1.8 mm. (a) Unmodified silane film, 24 h after defect formation; (b) silane film modified with CeO2, 48 h after defect formation; (c) silane film modified with TiO2, 24 h after defect formation

4.9 Schematic illustration of the reaction of CeO2 nanoparticles with reactive silanol groups

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The SVET measurements showed that the corrosion activity starting at induced defects was delayed in the presence of TiO2 or CeO2 nanoparticles (Figs. 4.7 and 4.8). However, for longer immersion times, only the CeO2 nanoparticles were highly effective in delaying the corrosion processes. The corrosion of steel substrates leads to the alkalinisation of cathodic sites. Under these conditions, the silica constituents of the silane coating start to decompose, because they are very unstable under alkaline conditions (pH > 9). Therefore, the silane layer enters a degradation process that is accelerated by the formation of a porous layer of expansive iron corrosion products. The growth of these corrosion products exerts interfacial stresses that promote the delamination of the silane coating. On the other hand, nanoparticles, especially CeO2, are highly stable under alkaline environments. Therefore, the nanoparticles can be released from the siloxane network during its decomposition and they may precipitate together with the corrosion products formed on the metallic surface. Most of the iron corrosion products, by themselves are not protective and corrosion proceeds very quickly. However, it was observed that these processes were significantly delayed in the presence of CeO2. It may be assumed that the nanoparticles enhance the protective performance of the corrosion products formed, slowing down the electrochemical processes. However, the mechanisms that explain the improved corrosion resistance of metallic substrates in the presence of TiO2 or CeO2 nanoparticles are not fully proven. It is reported in the literature that these nanoparticles have beneficial effects on corrosion resistance. The release of iron ions from steel substrates coated with polypyrrole/TiO2 composite films [25,26] decreased by about 50% compared to unmodified films. The beneficial effects of CeO2 are also reported in the literature. It was found that sulphidation of FeAl alloys was retarded because the nanoparticles inhibit iron diffusion and act as traps for the sulphur ions [26]. Cabot and Foissy also demonstrated that silica layers stabilised with CeO2 nanoparticles led to significant improvements in the corrosion resistance of zinc-coated steels [27]. Cerium oxides are also reported as having a pronounced stabilising effect on the passive state of steels and their corrosion resistance [28]. The results obtained in this work as well as results published in the literature [29–33], suggest that the precipitation of nanoparticles can allow the formation of more protective and stable layers of corrosion products that delay the corrosion activity. 4.5

Conclusions

Silane coatings obtained from bis-1,2-[triethoxysilylpropyl]tetrasulphide silane solutions containing TiO2 or CeO2 nanoparticles provide good barrier properties when applied on cold-rolled steel substrates. CeO2 addition resulted in better barrier properties. For the conditions tested in this work, TiO2 nanoparticles could delay corrosion activity during the early stages of immersion, but the effect vanished with increasing time of exposure to the aggressive solutions. Conversely, CeO2 nanoparticles induced a significant delay in the corrosion kinetics of steel substrates. Acknowledgments The authors acknowledge W. Trabelsi and P. Cecílio for their collaboration in the polarisation and SVET measurements, respectively.

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References 1. K. L. Mittal (ed.), Silanes and Other Coupling Agents, VSP, The Netherlands, 1992. 2. Y. Matsuda and H. Yasuda, Thin Solid Films, 118 (1984), 211. 3. M. Cabral, R. G. Duarte, M. F. Montemor and M. G. S. Ferreira, Corros. Sci., 54 (2005), 322. 4. A. Frignani, F. Zucchi, G. Trabanelli and V. Grassi, Corros. Sci., 48 (2006), 2258. 5. W. Trabelsi, P. Cecílio, M. G. S. Ferreira, K. Yasakau, M. L. Zheludkevich and M. F. Montemor, Prog. Org. Coat., 59 (2007), 214. 6. T. Van Schaftinghen, C. Le Pen, H. Terryn and F. Hörzenberger, Electrochim. Acta, 49 (2004), 2997. 7. M. F. Montemor and M. G. S. Ferreira, Electrochim. Acta, 52 (2007), 7486–7495. 8. M. F. Montemor, W. Trabelsi, M. Zheludkevich and M. G. S. Ferreira, Prog. Org. Coat., 57 (2006), 67. 9. M. G. S. Ferreira, R. G. Duarte, M. F. Montemor and A. M. P. Simões, Electrochim. Acta, 49 (2004), 2927. 10. L. E. M. Palomino, P. H. Suegama, I. V. Aoki, et al., Electrochim. Acta, 52 (2007), 7496– 7505. 11. V. Palanivel, Y. Huang and W. J. van Ooij, Prog. Org. Coat., 53 (2005), 153. 12. M. F. Montemor, A. M. Cabral, M. L. Zheludkevich and M. G. S. Ferreira, Surf. Coat. Technol., 200 (2006), 2875. 13. V. Palanivel, D. Zhu and W. J. van Ooij, Prog. Org. Coat., 47 (2003), 384–392. 14. 9. L. Liu, J.-M. Hu, J.-Q- Zhang and C.-N. Cao, Electrochim. Acta, 52 (2006), 538–545. 15. M. F. Montemor and M. G. S. Ferreira, Electrochim. Acta, 52 (2007), 6976–6987. 16. S. V. Mattigod, G. E. Fryxell, K. Alford, T. Gilmore, K. Parker, J. Serne and M. Engelhard, Environ. Sci. Technol., 39 (2005), 7306–7310. 17. F. Bauer, H.-J. Gläsel, E. Hartmann, H. Langguth and R. Hinterwaldner, Int. J. Adhes. Adhes., 24 (2004), 519–522. 18. W. Posthumus, P. C. M. M. Magusin, J. C. M. Brokken-Zijp, A. H. A. Tinnemans and R. van der Linde, J. Colloid Interf. Sci., 269 (2004), 109–116. 19. F. Grasset, N. Saito, D. Li, D. Park, I. Sakaguchi, N. Ohashi, H. Haneda, T. Roisnel, S. Mornet and E. Duguet, J. Alloys Comp., 360 (2003), 298–311. 20. E. Vassilev, B. Varimezova and K. Hadjiivanov, Anal. Chim. Acta, 336 (1996), 141–150. 21. M. Yan, W. Wei and N. Zuoren, J. Rare Earths, 25 (2007), 53–57. 22. X. Zhao, S. Ma, J. Hrbek and J. A. Rodriguez, Surf. Sci., 601 (2007), 2445–2452. 23. U. Diebold, Surf. Sci. Rep., 48 (2003), 53–229. 24. T.-S. Yang, C.-B. Shiu and M.-S. Wong, Surf. Sci., 548 (2004), 75–82. 25. D. M. Lenz, M. Delamar and C. A. Ferreira, J. Electroanal. Chem., 540 (2003), 35–44. 26. C. Xiao and W. Chen, Surf. Coat. Technol., 201 (2006), 3625–3632. 27. B. Cabot and A. Foissy, J. Mater. Sci., 33 (1998), 3945–3952. 28. E. Stoyanova, D. Nikolova, D. Stoychev, P. Stefanov and T. Marinova, Corros. Sci., 48 (2006), 4037–4052. 29. C. A. Ferreira, S. C. Domenech and P. C. Lacaze, J. Appl. Electrochem., 31 (2001), 49–56. 30. B. Mokshanatha Praveen, T. V. Venkatesha, Y. A. Naik and K. Prashantha, Synth. React. Inorg. Metal-Org. Nano-Metal Chem., 37 (2007), 461–465. 31. A. R. Phani, F. J. Gammel, T. Hack and H. Haefke, Mater. Corros., 56 (2005), 77–82. 32. M. F. Stroosnijder, M. J. Bennett, V. Guttmann, J. F. Norton and J. H. W. de Wit, Oxid. Metals, 35 (1991), 19–33. 33. P. Muhamed Ashrafa and S. M. A. Shibli, Electrochem. Commun., 9 (2007), 443–448.

5 Pyrrole-based silane primer for corrosion protection of commercial aluminium alloys

Monica Trueba and Stefano Trasatti Department of Physical Chemistry and Electrochemistry, University of Milan, Via Golgi 19, 20133 Milan, Italy [email protected]

5.1

Introduction

Many alternative technologies for the corrosion protection of aluminium alloys are under investigation, driven by the urgent need to replace chromate-based treatments due to the toxicity and carcinogenicity of hexavalent chromium ions. Among the more recent approaches, electronically conducting polymers (ECPs), namely polyaniline and polypyrrole, show many potential advantages as corrosion-inhibiting coatings due to both barrier action and anodic protection [1,2]. Nevertheless, processing difficulties due to their lack of solubility/fusibility and their low to moderate adhesion represent the main limitations to practical application. In parallel with ECPs, organosilane-based treatments, performed by immersing the metal in dilute alcohol- or water-based solutions for a short period, are also being intensively studied [3–5]. Effective coupling with both the metal substrate and organic topcoat can be obtained. Also, silane coatings are likely to be ‘passive’ since they act essentially as a physical barrier by hindering the penetration of aggressive species to the metallic substrate. Several modifications have been reported to make ECPs soluble and/or to improve their adhesion [6,7]. Doping silane coatings with small amounts of chemicals to provide inhibition properties has also been investigated [8,9]. Despite all of these studies, at present, no alternative has proven to be as reliable as chromate for corrosion protection. For successful replacement of chromate-based coatings, it is important to bear in mind that chromate treatments offer excellent corrosion resistance, including self-healing ability, as well as very good adhesion, and are easily applied. A pyrrole-based silane has been reported to promote polypyrrole (Ppy) adhesion on insulating substrates such as n-type silicon photoanodes [10], p-type silicon wafers [11] and glass fibres [12]. To the authors’ knowledge, this approach has not been used with reactive metals. If both ECPs and silane compounds are combined in a single coating, it should be possible to obtain a composite film with improved corrosion protection, chemical/mechanical stability and adhesion. In the end, simplification of the deposition and processing of ECPs would be feasible. This work describes a surface treatment on as-received commercial aluminium alloys with a pyrrole-based silane (SiPy). Coatings on 6082-T6, 5083-H111 and 2024T3 alloys were deposited following the classical steps for silane deposition, i.e. hydrolysis and condensation. The corrosion behaviour was assessed in chloride-containing 79

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80

solution by several electrochemical and chemical tests. Spectroscopic and microscopic techniques were used to characterise the coatings. 5.2 5.2.1

Experimental Materials and chemicals

Commercial wrought aluminium alloys AA6082-T6, AA5083-H111 and AA2024T3, with chemical compositions given in Table 5.1, were supplied by AVIOMETAL S.p.a. Plates with thicknesses of 1 to 1.5 mm (depending on the alloy) were cut into coupons with dimensions of 20 × 30 mm. All chemicals were of analytical reagent grade and were used as-received. 5.2.2

Surface treatment

No surface pre-treatment was carried out apart from ultrasonic cleaning in n-hexane, acetone and methanol, 15 min each. This surface condition was chosen to test an as-received surface in all substrates in accordance with the work of Jacob and colleagues [13], who reported that silane coupling agents are most effective on untreated aluminium alloys. The pyrrole-based silane (SiPy) molecule is schematically represented in Fig. 5.1. A SiPy solution was prepared at 4%v/v in methanol/water (95:5) with pH adjusted to 4 by adding acetic acid (10%). Methyltrimethoxysilane (Si) solution was prepared similarly. Preheated (120°C) metallic specimens were immersed for 3 min in SiPy or Si solutions (single immersion), dried in hot air and cured at 130–150°C for 2 h. Modified specimens were kept in a dessicator until used. 5.2.3

Characterisation techniques

X-ray photoelectron spectroscopy (XPS) analysis was carried out using an ESCA system (XI ASCII Surface Science Instruments) with an operating pressure of between 10–8 and 10–9 Torr, fitted with Al anode (1486.6 eV) and giving an energy resolution of 1 eV. FTIR measurements for pure and hydrolysed solutions of SiPy and Si were carried out with a Perkin-Elmer Spectrum 100 spectrophotometer, equipped with a constant horizontal angle reflectance accessory and a diamond crystal. Reflection-absorption IR (RAIR) spectra for the coated specimens were recorded on a Bio-Rad FTS-40 instrument, with a spectral resolution of 4 cm–1, and a scan number of 64. In all cases, the spectral range was 4000–400 cm–1.

Table 5.1 Chemical composition (wt%) of commercial wrought aluminium alloys (AA) Al alloy 6082 T6 5083 H111 2024 T3

Si

Fe

Cu

Mn

Mg

Zn

Ti

Cr

0.90 0.17 0.15

0.36 0.32 0.25

0.04 0.04 4.67

0.56 0.62 0.63

1.00 4.32 1.34

0.02 0.03 0.02

0.02 0.02 0.06

0.04 0.07 0.01

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Self-healing properties of new surface treatments

5.1 N-(3-(trimethoxysilyl)propyl)pyrrole (SiPy)

Scanning electron microscopy (SEM) images were obtained with a LEO 1430 microscope with a chamber pressure of 8 × 10–6 Torr and an accelerating voltage of 20 keV. 5.2.4

Corrosion tests

The working solution for these experiments was naturally aerated 0.6 M NaCl (pH 6.5 ± 0.2) prepared with distilled Milli-Q water and reagent grade NaCl (98%, Aldrich). Electrochemical studies were performed at room temperature in a singlecompartment O-ring cell [14] with a working (active) area of 1 cm2. A platinum sheet was used as a counterelectrode and with an external SCE, connected to the working compartment via a salt bridge containing the test solution and a Luggin capillary, as a reference electrode. Measurements were carried out with a microprocessorcontrolled potentiostat (Solartron 1286). Unless otherwise stated, all potentials are reported on the SCE scale. The pitting scans or single-cycle pitting curves at a rate (ν) of 10 mV/min were recorded after equilibration at open circuit for 10 min (teq). When the forward current reached about 5 × 10–3 A cm–2, the potential scan was reversed until complete repassivation had occurred. The open circuit potential (or free corrosion potential) for bare and modified substrates was monitored as a function of time for at least 15 h. Immersion tests were carried in test solution open to air at room temperature for a period of 7 days, in accordance with ASTM G31 recommendations [15]. 5.3

Results and discussion

5.3.1 5.3.1.1

Spectroscopic characterisation FTIR studies

Figures 5.2 and 5.3 illustrate high frequency (HW) and low frequency (LW) spectral regions, respectively, of SiPy in four forms: pure, hydrolysed (after 10 days), adsorbed layer and cured film on AA2024. Band assignments [16–23] for the observed infrared-active modes are compiled in Table 5.2; band features and positions very similar to those of SiPy/2024 were observed for SiPy on AA6082 and AA5083.

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5.2 FTIR spectra in the high frequency (HW) 3700–2700 cm–1 region for SiPy: (a) pure, (b) hydrolysed solution (after 10 days), (c) adsorbed layer on AA2024, (d) film on AA 2024

The high frequency region (HW) in the range 3700–2750 cm–1 (Fig. 5.2), mainly consists of ν(C–H) stretch vibrational modes of the pyrrole (Py) ring and alkane chain at ca. 3120–3100 cm–1 and 2975–2800 cm–1, respectively. As expected, hydrolysed and adsorbed layers (Fig. 5.2b,c) show a pronounced band between 3600 and 3100 cm–1, which corresponds to the ν(Si–OH) stretching mode from silanol groups in solution and those remaining in the organosilane layer on surface silanisation, respectively. For the former, the cleavage of Si–OCH3 bonds after hydrolysis is also manifested by the absence of symmetric stretchings νs(CH3) and νs(CH3,FR) of the methoxy groups at ca. 2840 and 2942 cm–1, showing up antisymmetric (νa) and symmetric (νs) stretching vibrations of CH2 in the propyl chain. A first indication that the siloxy network is completely condensed is given by the absence of ν(Si–OH) in the SiPy deposited on AA2024 (Fig. 5.2d). The complexity of the low frequency spectral region (LW) from 1750 to 450 cm–1, considerably increases when moving from pure to deposited SiPy (Fig. 5.3), in part

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Self-healing properties of new surface treatments

5.3 FTIR spectra in the low frequency (LW) 1750–450 cm–1 region for SiPy: (a) pure, (b) hydrolysed solution (after 10 days), (c) adsorbed layer on AA2024, (d) film on AA2024

because of the increased intensity of characteristic Py-ring bands. In the solution spectrum (Fig. 5.3b), the band at ca.880 cm–1 of ν(Si–OH), as well as the decrease in ν(Si–O–C) at 1075/804 cm–1 and of δ(CH3) at 772 cm–1, confirm SiPy hydrolysis. Characteristic Py-ring bands, i.e. in-plane vibrations, are detected at 1500, 1275, 1087 and 1064 cm–1. The latter two bands were obscured by the strong νa(Si–O–C) band at 1075 cm–1 in pure SiPy (Fig. 5.3a). To some extent, condensation in solution of some silanol groups occurs, according to the shoulders at 1117 and 1042 cm–1 that correspond to antisymmetric stretching vibrations of siloxane bonds νa(Si–O–Si). Partial crosslinking is justified by considering that the solution spectrum was recorded after 10 days of SiPy hydrolysis, in agreement with the well-known instability of silane solutions. In addition to the condensation of silanols, linking between a few Py units in the SiPy molecule is likely to occur, as suggested mainly by the weak features at ca. 800 and 1181 cm–1 [17], whose intensity increases in the SiPy adsorbed layer and film (Fig. 5.3c,d), which will be explained later on. An intense broad feature between 1150 and 1000 cm–1, splitting into several overlapping components, is observed for the SiPy adsorbed layer (Fig. 5.3c). Except for the two sharp central maxima at 1087 and 1064 cm–1, corresponding to in-plane C–H deformations of the Py ring δ(C–H)ip,ring, the shoulders at ca. 1110 cm–1 and

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Table 5.2 Infrared band assignments for SiPy as a pure liquid, hydrolysed solution, adsorbed layer and cured SiPy film on AA2024 Band frequency (cm–1) Pure liquid

Hydrolysed solution

Adsorbed on AA2024

Vibrational mode assignment Layer on AA2024

3678/3650 3590/3570 3123/3100 2968 2942 2914 2888/2875 2840

3600–3100 3126/3100 2962

3600–3100 3125/3099

2932 2882/2876 2845

2929 2874 2842 1719–1596 1494

3123/3099 2962 2933 2880/2875 2843

1500 1464 1445 1412 1358

1500 1461 1444 1406 1359

1278 n/o n/o 1189

1276 1240 1193

1312/1297 1275 1240 1195

1181

1180

1117 1086

1108 1085

1180 1139 1111 1086

1062 1043 970 913 880/852

1062 1038 970 914 880

1064 1040/1013 970 909 880

722 695 617 481w

801 721 697 615 483

1076 1067 n/o 970 n/o n/o 804 772 720 n/o 617 488

1441 1406 1354

1500 1462 1443 1409 1359 1343 1313/1296 1276 1240 1194

804 722 699 618 480

H2O H–O–Si ν(Si–OH) ν(Cα–H)ip,ring/ν(Cβ–H)ip,ring νa(CH3)/νa(CH2) νs(CH3,FR) νs(CH2,FR)/νa(CH2) νa(CH2)/νs(CH2) νs(CH3) H2O ν(C=C)ip,ring δ(CH2)scissor ν(C–N) δ(CH2)scissor δ(CH2) δ(CH2)gtma δ(CH2)Wxb ν(C–N)aliphatic δ(CH2)twist Py olig. doping-induced δ(CH3)rocking δ(C–H)ip,ring – B mode νa(Si–O–Si) νa(Si–O–Si) δ(C–H)ip,ring – T mode νa(Si–O–C) δ(C–H)ip,ring – T mode νa(Si–O–Si) ωa(C–H)oop,ring Py olig. doping-induced ν(Si–OH) νs(Si–O–C) δ(CH3) ωs(C–H)oop,ring – B mode δ(CH2)rocking ωs(C–H)oop,ring – T mode ring tor/ν(Si–O–C)?? νs(Si–O–Si)

a

gtm, end-gauche. Wx, coupled wag modes.

b

1138 cm–1 are assigned to νa(Si–O–Si). The presence of water and silanol groups is demonstrated by the bands between 1700 and 1600 cm–1 and at ca. 800 cm–1, respectively, which are no longer observed after heat treatment or curing (Fig. 5.3d). The spectrum for the film on AA2024 shown in Fig. 5.3d reflects a more pronounced broadening and splitting where the shoulder at 1110 cm–1 in the adsorbed

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Self-healing properties of new surface treatments

layer (Fig. 5.3c) is detected as a broad intense peak at nearly the same frequency. Moreover, two additional shoulders at ca. 1140 and 1013 cm–1 show up. These siloxane vibration features are consistent with the formation of very long linear chains and/or a mixture of cyclic segments. Upon condensation, Py-ring characteristic vibrations become sharper and other features appear that will be discussed below. Also, ν(C=C)ip,ring at 1500 cm–1 increases in intensity relative to the broad feature between 1150 and 1000 cm–1 and to the rocking deformation of the alkyl chain δ(CH2)rocking at ca. 720 cm–1, when compared to the SiPy solution and adsorbed layer spectra (Fig. 5.3b and c, respectively). Besides intense Py-ring characteristic vibrations, new bands are generated at ca. 1180 cm–1 and 800 cm–1, red shifted with respect to δ(C–H)ip,ring and ωs(C–H)oop,ring, respectively (see Table 5.2). Zerbi and colleagues have shown that these new features result from differences in the characteristic vibrations and intensities between the modes belonging to the two Py rings at either end of the chain (T – ‘end group’ modes) and those of pyrrole groups within the chain (B – ‘bulk unit’ modes) [17]. The frequency positions for T modes are blue shifted with respect to the B modes. In addition, as the number of pyrrole repeat units in the chain increases, the intensity of the B band in comparison to the T band for a given characteristic Py-ring vibration will increase. On this basis, the observed features are consistent with B mode generation due to the increase in length of the pyrrole chain. Moreover, the clear T and B modes at ca. 700 and 800 cm–1, respectively, for ωs(C–H)oop,ring, and the deformation mode of the two C–H bonds in position β−β’, indicate that the linking of the Py rings is substantially α−α’. In addition to B mode vibrations, the shoulder at ca. 1194 and the bands around 900 cm–1 (Fig. 5.3d) indicate some degree of doping of pyrrole oligomers. As is known, doping in polyconjugated materials generates at least two extra strong bands near 1195 and 910 cm–1 for polypyrrole. These results indicate that crosslinking of SiPy molecules occurs via both the condensation of silanol groups and α−α’ linking of Py rings, thus explaining the high degree of SiPy film cohesion observed by SEM (Fig. 5.4). The estimated thickness of the amorphous-like SiPy films was 10 μm for AA6082, 5 μm for AA5083 and 2 μm for AA2024. Additional information about the surface orientation of the SiPy film can be obtained from the LW region in Fig. 5.3 [21]. From the selection rules of RAIR spectroscopy, should a preferred orientation exist, there would be differences in the intensity ratios of a given pair of bands when moving from isotropic to reflection spectra. Two vibrational modes of the Py-ring whose transition dipole moments are orthogonal to each other, namely, ν(C=C)ip,ring at ca. 1500 cm–1 and ωs(C–H)oop,ring at ca. 700 cm–1, are very useful. This was already pointed out above and is particularly noticeable for the intensity ratios of these bands when comparing the SiPy solution and adsorbed layer spectra, indicating that Py-rings probably adopt a preferred orientation, mainly perpendicular to the surface, upon adsorption. The sharpness of the vibrations in the SiPy film are probably related to conformational rearrangements induced principally by the elimination of water on heat treatment. The conformational order of propyl chains can be inferred from analysis of the corresponding vibration modes [19,22]. The relative intensity ratios of methylene stretching vibrations νa(CH2) and νs(CH2) at ca. 2933 and 2878 cm–1, respectively, in the HW region (Fig. 5.2), do not significantly differ (2.00 in solution and 1.78 in reflection spectra), indicating that no change in the alkyl chain tilt angle has been produced upon SiPy deposition. In the LW spectral range (Fig. 5.3), the lowest frequency peak at ca. 722 cm–1, assigned to the CH2 rocking mode vibration, keeps its

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86

5.4 Cross-sectional SEM image of as prepared SiPy film on AA6082

sharp aspect and frequency, which suggests highly organised alkyl chains. The presence of kink defects will broaden and shift such vibration modes towards higher frequencies. The indication above can be further supported by a sharp feature observed at 1462 cm–1, related to the CH2 scissoring deformation mode. Unfortunately, the close vicinity of ν(C–N) at 1443 cm–1 does not allow determination of whether there is some band splitting. Furthermore, in the 1350–1295 cm–1 frequency range for the SiPy film (Fig. 5.3d), two well-resolved peaks are observed at 1313 and 1296 cm–1, respectively. These features are assigned to a progression of coupled CH2 wag modes whose appearance is evidence for a trans-conformational sequence in the alkyl chains. Also, the spacing of 17 cm–1 is in agreement with the theoretical value for a perfect all-trans chain, even though the appearance of a very weak, yet distinguishable shoulder at 1343 cm–1 is consistent with localised end-gauche (gtm) defects. The characterisation data presented so far suggest that crosslinking of SiPy molecules is produced by both the condensation of silanol groups and α−α’ linking of Py rings, the oligomers of which show some degree of doping. The network is highly crosslinked and free from gross defects, where the Py rings are preferentially oriented perpendicular to the surface and the propyl chains are mainly in transconformational sequence. Figure 5.5 illustrates the spectra obtained for methyltrimethoxysilane (Si) in three forms: pure, hydrolysed solution (after 10 days) and deposited on AA2024. Similar responses were obtained for Si film on AA6082 and AA5083. The number of vibrational bands significantly decreased in comparison to SiPy (Figs. 5.2 and 5.3), in

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Self-healing properties of new surface treatments

5.5 FTIR spectra for Si: (a) pure, (b) hydrolysed solution (after 10 days), (c) film on AA2024

accordance with the simple silane structure (Table 5.3) [18,22]. The broad features at ca. 1117 and 1017 cm–1 for Si film in Fig. 5.4c are assigned to νa(Si–O–Si) and are characteristic of linear polysiloxanes. The typical crystalline morphology of polymethylsiloxane films is shown in Fig. 5.6. Table 5.3 Infrared band assignments for Si pure liquid, its hydrolysed solution, and cured Si film on AA2024 Band frequency (cm–1) Pure liquid

Hydrolysed solution

Layer on AA2024 3586/3565

2970 2945 2841 1268 1191 1081 839 792

3600–3100 2971 2948 2847 1270 1191 1100 1045 879 777

2969 2850 1270 n/o 1117 1017 892 775

Vibrational mode assignment H–OSi ν(Si–OH) νa(CH3) νs(CH3,FR) νs(CH3) δ(CH3) δ(CH3)rocking νa(Si–O–Si) νa(Si–O–C) ν(Si–OH) νs(Si–O–C) δ(CH3)

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5.6 Cross-sectional SEM image of as prepared Si film on AA6082

5.3.1.2

XPS analysis

XPS spectra were obtained for SiPy film on AA6082, AA5083 and AA2024. Peaks for O1s, N1s, C1s and Si2p were clearly identified as shown for SiPy/AA2024 in Fig. 5.7. The results obtained from the high resolution analysis together with chemical assignments [20,24,25] are reported in Table 5.4. The complex C1s analysis is explained by the different carbon types in SiPy such as those of methylene groups and α-(or C–N) and β-carbons of the Py ring. Oxygencontaining groups can be attributed to surface oxidation of Py chains. But the absence of the carbonyl band at ca. 1700 cm–1 in RAIR spectra (Fig. 5.3d) may suggest that oxygen-containing groups include some contribution from acetate ions, expected to be involved in the doping of Py oligomers. In fact, for SiPy on AA6082 and AA5083, these are indicated by the highest binding energy peak (ca. 290 eV), in correspondence with the O1s peak at ca. 533 eV. The main component for the latter line region at ca. 532 eV, as well as the peak for Si2p at ca. 102 eV, are in agreement with binding energy positions for polysiloxanes, indicating their presence on the surface. The most interesting result is observed in the analysis of the N1s line, illustrated in Fig. 5.8. Two nitrogen species, namely, the imine-like (=N–) and positively charged nitrogen (N+) structures are obtained for SiPy/AA2024 (Fig. 5.8c), while only pyrrolic nitrogen (C–N) is detected in the case of SiPy on AA6082 and AA5083 (Fig. 5.8a,b). It has been reported that Al, Mg and Cu atoms can indirectly interact with the nitrogen of the pyrrole ring through the dopant of the polymer, giving =N– and N+ that cause disruption of the π electron configuration [25]. These interactions

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Self-healing properties of new surface treatments

5.7 XPS spectra of SiPy film on AA2024

Table 5.4 Binding energies (eV) and chemical assignments for High-Resolution XPS photopeaks of SiPy films on AA6082, AA5083 and AA2024 Photopeak

AA 6082

AA 5083

AA 2024

O1s

532.00 533.03

531.88 533.02

531.88

N1s

400.28

400.02

284.67 285.84

284.70 285.57

286.94 287.89 288.99 102.41

287.59 289.19 102.27

399.16

C1s

Si2p

401.04 284.58 285.67 286.60 288.02 102.18

Chemical assignment Si–O C–O/C=O Imine-like (=N–) C–N pyrrolic N+ C–H, Cβ Cα/C–N C=N C–O C=O O=C–O C–Si–O

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5.8 XPS high resolution analysis of N1s lines for SiPy film on: (a) AA6082, (b) AA5083, (c) AA2024

might have occurred between the AA2024 substrate and the SiPy film since, in addition to imine-like and N+ nitrogen constituting the main structure of Py-ring nitrogen, the peak for acetate groups was not observed. Moreover, the intrinsic oxidation state ratio [N+/N] is 0.27, which is comparable to that generally reported for as-synthesised or completely undoped polypyrrole films (≅0.25). During deposition with a solution pH of 4, destabilisation of Al oxide could readily occur, whose dissolution will produce Al3+ and Cu2+ (the main alloying element). These metallic ions might be trapped in the SiPy film and interact with the Py-ring nitrogen. The following question arises: why are the XPS results for SiPy on AA6082 and AA5083 not indicative of π electron configuration disruption, showing acetate groups and only pyrrolic N, if mainly Al3+ for the former and, both Al3+ and Mg2+, for AA5083, could be similarly produced and interact with Py-ring nitrogen? Indeed, this is not ruled out, although it is probably dependent on alloy reactivity or preferential dissolution. For AA6082 and AA5083, an aluminium rich surface would

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be exposed to the solution inside the pits, according to their typical pitting characteristics, which is accessible to silanol groups from the SiPy solution. On the other hand, Cu-rich aluminium alloys such as AA2024 are characterised by enrichment of the alloying element inside the pits, making the covalent bonding of silanol groups difficult over all the Al substrate [5]. The latter could form voids in polysiloxane chains, allowing more ‘free’ space for metallic ions to diffuse through the film and interact with Py-ring nitrogen. In other words, possible interaction of metallic ions released from AA6082 and AA5083 with Py-ring nitrogen should be mainly confined to the substrate/SiPy interface, not detectable by XPS. In the case of AA2024, the low affinity of the substrate surface for silanol group attachment (probably due to copper enrichment) favours penetration of the ions produced by substrate dissolution through the SiPy film. Additional results supporting the above considerations are discussed below. XPS-derived elemental stoichiometry ratios for Si/N and Si/O are summarised in Table 5.5. Carbon is excluded due to the great excess obtained, probably related to C contamination which is widely encountered in XPS studies. The surface composition of SiPy layers is reflected by Si/N ratios that are lower than the theoretical value for structure I, schematically depicted in Fig. 5.9 (left). If hydrogen-bonded silicon (or voids) are present to some extent, as indicated by the negative weak peaks at ca. 3590/3570 cm–1 in RAIR spectra for SiPy/AA2024 (Fig. 5.2d), schematically represented by structure II (Fig. 5.9, right), almost perfect agreement is obtained. Thus, the Si/N ratio tends to the theoretical value for structure I with decreasing alloy reactivity. A similar tendency is obtained for Si/O ratio, even though excess oxygen does not allow a good match between experimental and theoretical values. This suggests that the nature of the alloy surface will determine the extent of induced (chemical and/or physical) defects in the SiPy film, as well as its thickness. 5.3.1.3

Molecular picture of SiPy film

With the above surface spectroscopic data, it is difficult to understand how the SiPy film is built to have mainly Py-ring α–α’ linking, as well as to justify film thicknesses in the order of microns. A possible explanation is to consider that such linking is favoured in solution, where the SiPy molecules are ‘free’ to move and the α–α’ coupling could be catalysed by H+ ions (pH 4). Otherwise, steric hindrance and Py ring torsion will not support it. In fact, when comparing the transmission spectra of freshly prepared and 2-month aged SiPy solutions in Fig. 5.10, mainly analysing the B and T modes at ca. 700 and 800 cm–1, respectively, of the deformation mode ωs(C–H)oop,ring for the C–H bond in the β−β’ position, the increase in B band intensity is more

Table 5.5 XPS-derived elemental ratios for SiPy films on AA6082, AA5083 and AA2024 SiPy3 film

Si/N

Si/O

AA 2024 AA 5083 AA 6082 Theor. I Theor. II

2.11 2.63 2.76 3 2

0.52 0.58 0.72 1 0.66

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5.9 Theoretical structures I (left) and II (right)

5.10 FTIR spectra in the 3600–450 cm–1 region of SiPy hydrolysed solution freshly prepared (—) and aged for 2 months (---)

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5.11 Schematic representation of SiPy molecules in solution

pronounced with respect to that of the T mode. That is, further chain growth through α-α’ linking between Py oligomers is more favoured than the formation of new oligomers, in agreement with the Py polymerisation mechanism [26]. Under the same conditions, silanol groups are stable as indicated by the overlapping of the ν(Si–OH) stretching modes, i.e. the wide band between 3600 and 3100 cm–1 and that at ca.880 cm–1, in both fresh and old SiPy solutions. A diagram of the proposed conformational structure is shown in Fig. 5.11. Although highly idealised, it captures the essence of the structural features of SiPy molecules in solution. On adsorption, SiPy oligomeric molecules could be ‘frozen’ by the substrate surface mainly via Al–O–Si bonds, determining the preferred orientation of the SiPy film. That is, an apparent self-assembling of the layers upon fixing the molecules to the alloy substrate is produced, followed by a rearrangement of the SiPy molecules due to the condensation reaction of silanol groups, water/solvent elimination, and further polymerisation by Py rings. 5.3.2 5.3.2.1

Corrosion tests Pitting scans (PS)

Pitting scans are illustrated in Fig. 5.12 for bare, SiPy- and Si-treated AA6082 and AA5083 in naturally aerated 0.6 M NaCl solution. Curves for specimens coated by multiple deposition of the SiPy and Si films are included in the plot. SiPy treatments (Fig. 5.12a,b, left) produce a significant shift of the pitting forward-scans towards more noble potentials with respect to the bare substrate, this ennobling being more evident when multiple SiPy deposition is used. Also, the forward scan reveals that between the pitting potential and final breakdown, several zones of nearly constant current fluctuations are identified. That is, the forward scan is characterised by ‘semi-passivation’ zones followed by a rise in the current at nearly

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5.12 Pitting scans (ν = 10 mV/min) for SiPy (left) and Si (right) on: (a) AA6082 and (b) AA5083 in naturally aerated 0.6 M NaCl (pH 6.5 ± 0.2); (—) bare, (---) single, (—) multiple

constant potentials, giving an apparent ‘rising staircase’ shape. These features reflect the reformation of the protective oxide film within the defects or attacked areas, this repairing effect probably being induced or assisted by Ppy moieties in the SiPy film [27]. A remarkable hysteresis with a pronounced potential step characterises the

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reverse scan. Potential abruptness during the reverse scan for bare alloys has been called the pit transition potential (Eptp) [28], which represents an intrinsic property of the alloy and is associated with repassivation events. The similarity of Eptp for bare and coated AA6083 and AA5083, at ca. –740 mV, indicates that the first stages of surface repassivation are determined by the underlying substrate, where apparently the SiPy film is not directly involved. The latter apparently limits the complete repassivation of the surface. The pitting scans for Si-treated AA6082 and AA5083 are presented in the right part of Fig. 5.12a,b. Single immersion treatment provides some barrier protection as indicated by the lower passive currents with respect to the bare alloys. However, the multiple approach is better than the former with lower currents and wider passive regions. These differences are attributed to defects hidden by layer-on-layer deposition. The hysteresis on the reverse scan has the same features as those discussed above, irrespective of the film thickness. By comparing the PS of SiPy- and Si-modified AA6082 and AA5083, on the left and right parts of Fig. 5.12, respectively, it is observed that Si multiple treatment behaves similarly to a one-step SiPy deposition. The shift of the pitting onset for the latter is as much as 200 mV higher than that obtained for Si. Furthermore, after initial breakdown, the potential–current relation is predominantly linear. The notably different pitting scan behaviour of AA2024 treated with SiPy and Si is illustrated in Fig. 5.13. No passive regions are obtained for any treatment condition. The corrosion potential of the SiPy-coated alloy is shifted towards the noble direction by ca. 100–150 mV with respect to the response of bare metal (Fig. 5.13, left), while the multiple Si-treated substrate retards pit propagation to some extent (Fig. 5.13, right). Upon reversing the scan, the hysteresis closely follows that of the bare alloy but with a less pronounced potential step (Eptp).

5.13 Pitting scans (ν = 10 mV/min) for SiPy (left) and Si (right) on AA2024 in naturally aerated 0.6 M NaCl (pH 6.5 ± 0.2); (—) bare, (---) single, (—) multiple

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The inferior performance of coated AA2024 is associated with the less efficient bonding of silanol groups to the metal surface as a result of the alloy composition and pitting characteristics, thus suggesting lower surface affinity for the studied silane-based treatments with respect to AA6082 and AA5083. 5.3.2.2

Potential transient measurements (Et)

The free corrosion potential (Ecorr) as a function of time for SiPy- and Si-coated alloys (by single immersion), is illustrated in Figs. 5.14 and 5.15, respectively. Uncoated AA6082 and AA5083 display some degree of protection towards anodic dissolution of surface oxide before pitting starts, as can be seen in either Fig. 5.14a,b or Fig. 5.15a,b, characteristic of these bare alloys. Flat potential transients, initially positive that pass through a maximum and stabilise at ca. –850 mV until the end of the experiment, characterise both SiPy-treated AA6082 and AA5083 (Fig. 5.14a,b). This trend is indicative of initial polymeric film conformational relaxation [29], followed by O2 reduction at the film surface [30]. This reflects very good film performance for corrosion protection of these alloys, in agreement with the results obtained from the pitting scans (Fig. 5.12, left). Much longer times are needed for film breakdown to occur under this test condition. In the case of SiPy/AA2024 (Fig. 5.14c), the potential–time curve exhibits many transients according to the strong Ecorr fluctuations, which are mainly anodic with respect to the bare substrate. These potential fluctuations tend to lower and lower values of both ‘amplitude’ and ‘intensity’ within the first 3 h in NaCl solution, then showing some occasional transients to the bare substrate potential until the end-time of the experiment. Although this result indicates that SiPy film performance on AA2024 is worse with respect to that on AA6082 and AA5083, some repairing effect is inferred, which could not be deduced from the pitting scan experiments (Fig. 5.13, left). In any case, surface treatment with polymethylsiloxane shows significant protection, as illustrated in Fig. 5.15. The negative potential transient for Si/AA5083 (Fig. 5.15b), though not as low as for uncoated substrate (at ca. –1350 mV), indicates some degree of protection towards anodic dissolution of surface oxide. Nevertheless, the coating fails after several hours, reproducing the rest potential of the bare alloy. In the case of Si/AA6082 and Si/AA2024, the response closely follows that of the bare substrates over the duration of the experiment (Fig. 5.15a,c). The morphology of the cross-section for Si- and SiPy-coated AA5083 and AA2024 at the end of Et experiments is illustrated in Fig. 5.16. The pitting promoting chloride ions can easily attack the crystalline Si film, thus inducing significant corrosion with subsequent film deterioration (Fig. 5.16a,b), while the amorphous-like SiPy films with a high degree of cohesion and strongly adhered to the metal substrate provide significant corrosion protection (Fig. 5.16c,d). Isolated corrosion products and cracks are observed on SiPy/AA5083 and, even if an irregular film is obtained on AA2024, the adhesion is still good and the film failure occurs mainly at the SiPy surface rather than at the SiPy/metal interface, as indicated by the localisation of the corrosion products. The same can also be inferred for SiPy/AA5083. These results can be interpreted by considering that under free corrosion conditions, Ppy moieties in the SiPy film promote cathodic oxygen reduction at the SiPy/ solution interface distancing this reaction from the metal/film interface and, hence, avoiding film disbondment [30]. Moreover, the metal ions produced by aluminium dissolution beneath the coating sustained by oxygen reduction can be transported

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5.14 Potential transient curves for bare and SiPy-coated Al alloys: (a) AA6082, (b) AA5083, (c) AA2024; in naturally aerated 0.6 M NaCl (pH 6.5 ± 0.2); (—) bare, (·····) SiPy

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5.15 Potential transient curves for bare and Si-coated Al alloys: (a) AA6082, (b) AA5083, (c) AA2024, in naturally aerated 0.6 M NaCl (pH 6.5 ± 0.2); (—) bare, (·····) SiPy

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5.16 Cross-sectional SEM images for: (a) Si/AA5083, (b) Si/AA2024, (c) SiPy/ AA5083, (d) SiPy/AA2024, after Et experiments

through the film to the electrolyte; that occurs mainly over the SiPy surface for AA5083 and at the metal in the defects for AA2024. 5.3.2.3

Immersion tests

These experiments were performed on SiPy-coated alloys. Visual examination of bare and treated alloys after immersion for 7 days in 0.6 M NaCl (pH 6.5 ± 0.2) is illustrated in Fig. 5.17. High levels of long-term corrosion protection are observed for coated AA6082 and AA5083, virtually free of any corrosion product (Fig. 5.17a,b, right). This is particularly evident for the latter with a bright metal-like appearance, while the uncoated substrate shows characteristic opaqueness due to the presence of corrosion products. In the case of SiPy/AA2024 (Fig. 5.17c, right), significant localised corrosion mainly at the edges is observed on coated specimens, even though a still shiny metal-like appearance dominates the central area of the sample. The final pH value of the test solutions supports the above observations. A pH value close to 6.5 (the initial pH) was obtained for SiPy/AA6082 and SiPy/AA5083, while significant solution alkalinisation was detected in the case of SiPy/AA2024. The results of SEM examination of these surfaces are shown in Figs. 5.18 and 5.19, which further confirm the high degree of corrosion prevention by direct-to-metal pyrrole-based silane deposition with a single immersion for a short time. The typical crystallographic pitting for bare AA6082 (Fig. 5.18a) and roughly rounded grains of corrosion products on AA5083 (Fig. 5.18c), are not observed for the coated alloys (Fig. 5.18b,d). In addition, the sparsely found areas of some undercoating corrosion

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5.17 Bare (left) and SiPy (right) films on: (a) AA6082, (b) AA5083 and (c) AA2024 after immersion for 7 days in naturally aerated 0.6 M NaCl solution (pH 6.5 ± 0.2)

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5.18 SEM micrographs for bare and SiPy-treated alloys: (a) bare AA6082, (b) SiPy/AA6082, (c) bare AA5083, (d) and (e) SiPy/AA5083, after immersion for 7 days in naturally aerated 0.6 M NaCl solution (pH 6.5 ± 0.2)

on SiPy/AA5083, probably at the defects and/or due to porosity, are characterised by a eutectic-like Al–Si solidification [31] as illustrated in Fig. 5.18e. This suggests that the SiPy surface is highly flat, which is a prerequisite for observation of the Al–Si eutectic phenomenon. Significant undercoating corrosion is obtained for SiPy/AA2024, as demonstrated in Fig. 5.19a. However, quite efficient protection is still given by SiPy as indicated by the repair of defects (Fig. 5.19b), due to Al oxide regeneration probably assisted by the Ppy moieties in the SiPy film. An image of non-protected AA2024 at the same magnification is included (Fig. 5.19c). Accordingly, the inferior performance in the case of AA2024 is mainly related to the low affinity between the silane and the Cu-rich surface of the substrate, affecting film adhesion and promoting irregularities in the film thickness as a result of voids (or defects) in the SiPy film.

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5.19 SEM micrographs for bare and SiPy-treated AA2024: (a) and (b) SiPy/ AA2024, (c) bare AA2024, after immersion for 7 days in naturally aerated 0.6 M NaCl solution (pH 6.5 ± 0.2)

103 5.4

Self-healing properties of new surface treatments Conclusions

The deposition of a pyrrole-based silane (SiPy) film on aluminium alloys following the classical steps for the deposition of silanes, i.e. immersion in the SiPy solution and curing, results in a composite coating that includes both polysiloxane and polypyrrole units. The crosslinking of SiPy molecules is produced by both the condensation of silanol groups and α–α’ linking of pyrrole (Py) rings with some degree of doping. Pyrrole rings are preferentially oriented perpendicularly to the surface, and propylic chains mainly in trans-conformational sequence. The highly crosslinked network supports the well-packed highly coherent film morphology. To justify film thicknesses of the order of microns, from 10 μm for AA6082 to 2 μm for AA2024, as well as the prevailing α–α’ coupling of Py rings in the SiPy molecule, together with the experimental observation of readily occurring oligomerisation of Py rings in the hydrolysed solution, it is considered that the rotation of molecules is favoured in the latter, determining the high degree of order of the SiPy film upon adsorption on the metal surface. Impressive corrosion protection is given by the SiPy films in comparison to simple polysiloxane (polymethylsiloxane). SiPy corrosion prevention is better when multilayer deposition is carried out on a pre-heated metal substrate with longer immersion times and high-temperature curing. Polysiloxane linkages in the SiPy layer contribute largely to the improved barrier protection, while Ppy units may assist oxide film formation at the defects and/or alleviate the cathodic disbondment by promoting oxygen reduction on the SiPy film surface. References 1. (a) D. E. Tallman, G. Spinks, A. Dominis and G. G. Wallace, J. Solid State Electrochem., 6 (2002), 73–84; (b) G. Spinks, G. G. Wallace and D. E. Tallman, ibid, 85–100. 2. P. Zarras, N. Anderson, C. Webber, D. J. Irvin, J. A. Irvin, A. Guenthner and J. D. Stenger-Smith, Radiat. Phys. Chem., 68 (2003), 387–94. 3. M. A. Petrunin, A. P. Nazarov and Y. N. Mikhailovski, J. Electrochem. Soc., 143 (1996), 251–7. 4. (a) B. B. Johnsen, K. Olafsen, A. Stori and K. Vinje, J. Adhes. Sci. Technol., 16 (2002), 1931–48; (b) B. B. Johnsen, K. Olafsen, A. Stori and K. Vinje, ibid, 17 (2003), 1283–98. 5. (a) D. Zhu and W. J. Van Ooij, Corros. Sci., 45 (2003), 2163–75; (b) D. Zhu and W. J. Van Ooij, ibid, 2177–97. 6. V. J. Gelling, M. M. Wiest, D. E. Tallman, G. P. Bierwagen and G. G. Wallace, Prog. Org. Coat., 43 (2001), 149–57. 7. E. Hür, G. Bereket and Y. Ğahin, Mater. Chem. Phys., 100 (2006), 19–25. 8. L. E. M. Palomino, P. H. Suegama, I. V. Aoki, Z. Paszti and H. G. De Melo, Electrochim. Acta, 52 (2007), 7496–7505. 9. J. H. Osborne, K. Y. Blohowiak, S. R. Taylor, C. Hunter, G. Bierwagon, B. Carlson, D. Bernard and M. S. Donley, Prog. Org. Coat., 41 (2001), 217–25. 10. R. A. Simon, A. J. Ricco and M. S. Wrighton, J. Am. Chem. Soc., 104 (1982), 2031–4. 11. C. G. Wu and C. Y. Chen, J. Mater. Chem., 7 (1997), 1409–13. 12. F. Faverolle, A. J. Attias, B. Bloch, P. Audebert and C. P. Andrieux, Chem. Mater., 10 (1998), 740–52. 13. H. Woo, P. J. Reucroft and R. J. Jacob, J. Adhes. Sci. Technol., 7 (1993), 681–97. 14. J. N. Murray, Prog. Org. Coat., 30 (1997), 225–33. 15. Standard Practice for Laboratory Immersion Corrosion Testing of Metals, G31–72 (99), ASTM, West Conshohocken, PA, USA.

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16. B. A. Ashby, J. Organomet. Chem., 5 (1966), 405–12. 17. (a) B. Tian and G. Zerbi, J. Chem. Phys., 92 (1990), 3886–91; (b) B. Tian and G. Zerbi, ibid, 3892–98; (c) G. Zerbi, M. Veronelli, S. Martina, A. D. Schlüter and G. Wegner, ibid, 100 (1994), 978–84. 18. W. R. Thompson, M. Cai, M. Ho and J. E. Pemberton, Langmuir, 13 (1997), 2291–2302. 19. A. N. Parikh, M. A. Schivley, E. Koo, K. Seshadri, D. Aurentz, K. Mueller and D. L. Allara, J. Am. Chem. Soc., 119 (1997), 3135–43. 20. H. Kariis, E. Smela, K. Uvdal, M. Wirde, U. Gelius and B. Liedberg, J. Phys. Chem. B, 102 (1998), 6529–38. 21. R. L. McCarley and R. J. Willicut, J. Am. Chem. Soc., 120 (1998), 9296–304. 22. M. Cai, M. Ho and J. E. Pemberton, Langmuir, 16 (2000), 3446–53. 23. D. L. Elmore, D. B. Chase, Y. Liu and J. F. Rabolt, Vib. Spectrosc. 34 (2004), 37–45. 24. M. Omastova, K. Boukerma, M. M. Chehimi and M. Trchova, Mater. Res. Bull., 40 (2005), 749–65. 25. (a) V. W. L. Lim, S. Li, E. T. Kang, K. G. Neoh and K. L. Tan, Synth. Met., 106 (1999), 1–11; (b) V. W. L. Lim, E. T. Kang and K. G. Neoh, Macromol. Chem. Phys., 202 (2001), 2824–31; (c) V. W. L. Lim, E. T. Kang, K. G. Neoh and W. Huang, J. Vacuum Sci. Technol. A, 19 (2001), 2680–8. 26. A. Malinauskas, Polymer, 42 (2001), 3957–72. 27. T. D. Nguyen, M. Keddam and H. Takenouti, Electrochem. Solid State Lett., 6 (2003), B25–8. 28. M. Yasuda, F. Weinberg and D. Tromans, J. Electrochem. Soc., 137 (1990), 3708–15. 29. E. Krivan, C. Visy and J. Kankare, J. Phys. Chem. B, 107 (2003), 1302–8. 30. (a) P. J. Kinlen, D. C. Silverman and C. R. Jeffreys, Synth. Met, 85 (1997), 1327–32; (b) T. Schauer, A. Joos, L. Dulog and C. D. Eisenbach, Prog. Org. Coat., 33 (1998), 20–7. 31. K. Nogita and A. K. Dahle, Mater. Char., 46 (2001), 305–10.

6 Sol–gel derived hybrid materials as functional coatings for metal surfaces Johanna Kron, Karl-Joachim Deichmann and Klaus Rose Fraunhofer ISC, Neunerplatz 2, 97082 Wuerzburg, Germany [email protected]

6.1

Introduction

There are many reasons for protecting metal surfaces with coatings. The fight against corrosion is the most important aim from the economic point of view [1,2]. Here, the focus of interest is not only on the loss of material but also on the decrease and/or slowdown of corrosion, or – if damage has already occurred – on the restoration of the corroded metal parts. The protection of highly glossy metal surfaces which are often very sensitive to soiling and scratching [3–5] is accomplished by applying transparent coatings. Moreover, the use of coloured layers opens the pathway to decorative surface finishings. At the Fraunhofer-Institut für Silicatforschung, hybrid sol–gel materials accessible by chemical nanotechnology have been developed. In recent years, these materials have proved to be suitable candidates to tailor or alter metal substrates since these materials provide particular advantages [6–8]. For example, hybrid polymer sol–gel materials are potential substitutes for hexavalent chromium-based surface treatments. 6.2

Sol–gel technology – hybrid polymers

Sol–gel technology represents an important domain of chemical nanotechnology and is a key and interdisciplinary technology of the 21st century (world market in 2006: 800 million Euros [9]). With the help of sol–gel chemistry, the characteristics of non-metallic, inorganic and hybrid materials may be successfully implemented in a single material via nanoscaled intermediates. In principle, the sol–gel process is suited for manufacturing varied products such as powders, fibres, gels (aero- and xero-gels), bulk materials (e.g. components or devices) and layers (Fig. 6.1). In industry, the sol–gel process is used mainly in the field of surface technology for the production of functional layers. Hybrid polymer materials are synthesised by the sol–gel process through the controlled hydrolysis and condensation of organically modified Si-alkoxides. Cocondensation with other metal alkoxides (Ti-, Zr-, Al-alkoxides) is also possible (Fig. 6.2). On the one hand, functional organic groups are introduced to modify the inorganic network. On the other hand, polymerisable organic groups which are fixed to the inorganic network react with each other in thermal or UV-initiated processes to form additional organic polymeric structures (Fig. 6.3). The sol–gel process leads to clear sols which can be applied by conventional coating techniques such as dipping, spraying, flow-coating, roll-coating, screen printing, etc. The molecular precursors and processing conditions used in the sol–gel process 105

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6.1 Process techniques and materials via the sol–gel method

6.2 Hydrolysis and condensation reactions within the sol–gel process (R = Me, Et, Pr, etc; M = Si, Ti, Zr, Al, etc.)

are highly variable. That is why this method opens many pathways for material synthesis. The metal alkoxides build up inorganic structures and contribute to properties such as abrasion resistance, outdoor stability and adhesion to metal and metal oxide surfaces, whereas the organic moieties are responsible for the processability at low

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6.3 Organic crosslinking reactions to build up organic polymeric structures. R =functional organic groups like: –CnH2n+1, C6H5, –CnH2nCmF2m+1, –CnH2nNH2, etc.; X = polymerisable functions like: epoxy-, (meth)acrylic-, vinyl groups, etc

temperatures. Moreover, they contribute to the variability of the surface polarities, as well as to the flexibility and plasticity of the coated substrates. In addition, the properties of the hybrid sol–gel materials may be varied over a wide range by altering the reaction parameters such as temperature, time and pH value. The synthesis, upscaling and storage of the hybrid sol–gel materials can be controlled by applying different spectroscopic (e.g. RAMAN-, IR- and NMR-spectroscopy) and analytical (e.g. SEC, XPS) methods [10]. As shown in Fig. 6.4, the synthesis of such a hybrid nanocomposite coating is controlled by RAMAN spectroscopy. The epoxy-groups linked to the silanes which are contained in the system (band at 1258 cm–1 [10]) are still present after the hydrolysis of the alkoxy silane groups, while

6.4 RAMAN-spectra of a hybrid sol (main constituents: epoxy functionalised trimethoxy-silane compounds) (a) before hydrolysis (0.1 kg batch), (b) after hydrolysis (0.1 kg batch), (c) after hydrolysis (10 kg batch), (d) after hydrolysis and after 3 months of storage at –18°C (10 kg batch)

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the hydrolysis is indicated by the disappearance of the Si(OMe)3 band at 645 cm–1 [10]). These epoxy-groups are required for setting up organic polymer structures in the next step. Furthermore, the RAMAN-spectra indicate that the hybrid materials can be synthesised by the sol–gel process in different batch sizes which, nevertheless, result in materials with identical qualities. In addition, Fig. 6.4 presents an unchanged RAMAN-spectrum of a sol which was stored for 3 months at –18°C. The epoxy groups necessary for the organic crosslinking reactions are still present. 6.3

An alternative to chromating: hybrid sol–gel (ORMOCER®) based pretreatment of aluminium alloys

As a result of features such as excellent strength-to-weight ratio, easy processability, moderate corrosion resistance as well as physiological and environmental compliance, aluminium and its alloys are essential and valuable materials for a wide range of applications. Their good corrosion resistance is the result of an ‘in-situ’ formed, thin tenacious oxide film on the surface. Nevertheless, a chromating process followed by an organic finish is a common way to protect aluminium alloys for outdoor applications (e.g. in architecture) against corrosion. Figure 6.5 displays two examples of how to protect aluminium substrates with organic finishes. Despite the well-known efficiency of chromate pre-treatments to improve paint adhesion and corrosion resistance, the high toxicity and carcinogenic nature of Cr6+ compounds lead to environmental health risks and thus exert continuous pressure to develop suitable alternative corrosion protection systems. The high toxicity and resulting environmental hazards of Cr6+-containing compounds has caused environmental legislation to impose severe restrictions on chromate processes: The US Environmental Protection Agency (EPA) has listed chromium among the top toxic substances [11], and the EU directives ‘end-of-life vehicles’ [12] and ‘restriction of the use of certain hazardous substances in electrical and electronic equipment’ [13] have set new standards for the stepwise reduction in the application of procedures containing hexavalent chromium compounds. As a result of national legislation and because of market demands, process plants using

6.5 Examples of the layer structures on coated aluminium for outdoor applications

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chromate treatments will have to be shut down progressively over the next few years, if they are not able to establish alternative, environmentally friendly processes. Some alternative chromate free pretreatment methods for aluminium surfaces have already been described [14–16]. Commercially available products mainly consist of titanium and zirconium compounds. Hybrid polymer sol–gel materials are potential substitutes for hexavalent chromium-based surface treatments as well. Due to the chemical characteristics, in particular, the presence of hydroxy and alkoxy groups, the hybrid materials are qualified to coat metal as well as metal oxide surfaces. These groups can react with OH-groups on the surfaces of both metals and metal oxides. Water and alcohols are eliminated, while bonds between the hybrid polymer and the metal surface are created (Fig. 6.6), thus leading to good durable adhesion of the layers to the metal substrates [17,18]. Similar to silane adhesion promoters, the hybrid sol–gel materials can also link to organic polymer paint systems (Fig. 6.7). A major goal of research projects has been the adjustment of the chemical composition of the hybrid material to optimise the linkage to aluminium surfaces and organic finishes. Suitable conditions for an intermediate drying step for hybrid polymers and the final hardening of the layers can be determined by solid state carbon nuclear magnetic resonance (13C-NMR) spectroscopy. Figure 6.8 shows four solid state 13C-NMR-spectra for an epoxy functionalised hybrid polymer which was treated under different curing conditions. As a result of these investigations, the

6.6 Examples for creating bonds between sol–gel layer and metal surface

6.7 Examples for creating bonds between sol–gel film and organic layers

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6.8 Solid state 13C-NMR-spectra of a hybrid sol–gel material (main components: epoxy functionalised silane) cured at different conditions (a) 100°C, 10 min; (b) 160°C, 120 min; (c) 180°C, 40 min; (d) 220°C, 20 min

intermediate drying step of the hybrid polymers was accomplished at temperatures of approximately 100°C, since, under these conditions, epoxy groups remain unchanged within the system (respective NMR-resonances at 44.5 and 51.7 ppm [19]). At curing temperatures of about 160°C, the epoxy groups start to react within cross-linking reactions and the hybrid polymer can bond to the subsequently applied organic layer. A variety of hybrid sol–gel surface treatments applied to aluminium alloys such as AlMg1, AlMgSi1 (coil substrates) or AlMgSi0.7 (profiles) have been investigated. Together with the sol–gel pre-treatment, a conventional wet lacquer system consisting of a 2K-EP primer and a 2K-PUR topcoat as well as PES powder coatings have been investigated. In the case of the wet lacquer finish, an additional objective was to replace the chromating process as well as the primer layer. For the wet lacquer and powder coatings, sol–gel pre-treatments were identified which led to nearly perfect, durable adhesion (Gt 0/0) in the cross-cut test (according to ISO 2409) even after exposure to wet climates in the boiling test (according to DIN 58196-2/C60) or in the humidity test (condensed water test according to DIN EN ISO 6270-2). Most of the samples were stable in the bend test described in DIN EN ISO 1519; 5 mm thorn). The best coatings, which were based on epoxy functionalised silane compounds and aluminium alkoxides, resisted even more extreme testing conditions, e.g. 1000 h ESS test (acetic acid salt spray test; ISO 9227; Fig. 6.9). The test results on wet-lacquered sheets using an organic epoxy-based primer were equivalent to those without any primer. This means that both the chromating and the application of a primer were successfully replaced by a single process (Table 6.1).

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6.9 Aluminium sheets (alloy: AlMg1) covered with different organic finishes after 1000 h of ESS test, the lower parts of the sheets have been pretreated with hybrid sol–gel materials; left: powder coating (PES/Primid); centre: wet coating system (2K-EP primer and 2K-PUR top coat); right: 2K-PUR topcoat Table 6.1 Processing conditions and properties of hybrid sols and coated aluminium substrates (AlMg1) Properties

Range/valuation

Method

Viscosity Solids content Density Flash-point Dip-coating parameters Topcoat

2.5–4.5 mm2/s 2–5% ca. 1000 kg/m3 >323 K Dipping and drawing velocity: 100 mm/min; exposure time: 3 s Wet coating Wet coating Powder coating system: system: system: 2K-EP primer/ 2K-PUR top PES/Primid2K-PUR top coat coat powder coating <1 μm <1 μm <1 μm

Ubbelohde DIN EN ISO 3251 Pyknometer DIN EN ISO 3680

Profilometer

90 μm

60 μm

81 μm

Profilometer

Gt 0 Gt 0

Gt 0 Gt 0

Gt 0 Gt 0

ISO 2409 DIN 58196-2/C60

Gt 0

Gt 0

Gt 0

DIN EN ISO 6270-2

Stable

Stable

Stable

Stable

Stable

Stable

ISO 1519; 5 mm thorn ISO 9227; 1000 h

Coating thickness: sol–gel Coating thickness: topcoat Adhesion Adhesion after boiling test Adhesion after condensed water test Stability in the bend test Stability in the acetic acid salt spray test

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Nanocomposite protective layers for different metal surfaces

In layer thicknesses smaller than 1 μm, nanocomposite layers have proved to be potential adhesion promoters and suitable substitutes for hexavalent chromiumbased surface treatments. In layer thicknesses of more than 5 μm, they are expected to protect different metal surfaces against corrosion even as a single layer, due to the high variability of their chemical composition. The advantage of colourless, transparent nanocomposite layers is that the optical impression of the particular metal surface is retained. Additionally, highly glossy metal surfaces which are very sensitive to scratching, soiling and fingerprints can be equipped with a wear-resistant antiadhesive layer. Moreover, decorative effects are feasible with the application of transparent or opaque colour layers. Initially, nanocompound formulations were selected for synthesis and application which exhibited good adhesion to metal surfaces (see above). Furthermore, chemical approaches which were successful in the design of functional coatings for glass surfaces have been taken into consideration [20,21]. For the coating of glass surfaces, a composition starting from phenyl-tri(methoxy)silane, 3-glycidyloxypropyl-tri (methoxy)silane, 3-aminopropyl-tri(ethoxy)silane and aluminium-tri(sec-butoxide) had, amongst others, proved of value. By adding hydrophobic epoxide resins, which are soluble in the reaction medium and which cross-link with the epoxy-functionalised organo(alkoxy)silane, a considerable increase in the resistance against aqueous media could be achieved. Table 6.2 summarises the characteristics of the aforementioned hybrid polymeric sol and the resulting layer properties on different metal substrates. When stored at –18°C, the sols remain usable for 6 months. The layer properties of the sols stored

Table 6.2 Processing conditions and properties of thermally curing hybrid polymeric sols and coatings for metal surfaces Properties

Range/valuation

Method

Viscosity Solids content Density Flash-point Spray-coating parameters

9.5 . . . 19.8 mm2/s 35 . . . 55% ca. 1 kg/m3 301 K . . . 338 K Nozzle diameter 0.2–1.4 mm; pressure 1.5–3 bar Dipping and drawing velocity: 100 mm/min; exposure time: 3 s 433 K/2 h . . . 473 K/10 min 5 . . . 20 μm Gt 0 Gt 0 Gt 0 Gt 0 Gt 0 Gt 0 0.3 . . . 2.3% Stable (on top of all metal surfaces)

Ubbelohde DIN EN ISO 3251 Pyknometer DIN EN ISO 3680

Dip-coating parameters Curing conditions Layer thickness Adhesion to stainless steel Adhesion to copper Adhesion to brass Adhesion to aluminium Adhesion to silver Adhesion to gold-plated silver Abrasion resistance Adhesion after condensed water test

Profilometer ISO 2409 ISO 2409 ISO 2409 ISO 2409 ISO 2409 ISO 2409 DIN 52347 DIN EN ISO 6270-2

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6.10 Copper and aluminium substrates partially coated with hybrid sol–gel layers, after 14 days in the condensed water test (DIN EN ISO 6270-2); on top of the copper substrate, a cross cut test (DIN EN ISO 2409) and a taber abraser test (DIN 52347) is visible in addition

for different periods of time were comparable to those of freshly produced sols. Figure 6.10 demonstrates the protective function of a nanocomposite coating for aluminium surfaces when stored under a wet climate in the condensed water test (DIN EN ISO 6270-2). The double function of mechanical protection (abrasion resistance in the taber abraser test according to DIN 52347) and sufficient corrosion protection (chemical resistance according to DIN EN ISO 6270-2) of a well adherent hybrid polymer layer (Gt 0 according to DIN EN ISO 2409) can be seen on a copper surface. Another task in the further development of these thermally curing hybrid coatings was the implementation of a low surface energy which, in turn, makes them resistant to fingerprints and other contamination. The development of low surface energy sol–gel coatings has been an important topic in numerous investigations in recent years [22,23]. For other hybrid materials, it is known that the incorporation of long chain perfluoroalkyl silane compounds leads to low energy surfaces. The epoxysilane-based sol–gel system described above could also be equipped with perfluoroalkyl silanes, resulting in clear, smooth, homogeneous coatings by both the dip and spray coating techniques. The long chain perfluoroalkyl silane compounds are enriched in the upper boundary surface, as can be confirmed by XPS-investigations (Fig. 6.11) – irrespective of the quantities of perfluoroalkyl silanes used. Thus only very small quantities (up to 1 mol-%) of perfluoroalkyl silanes are necessary in the lacquer to generate low surface energy sol–gel layers. Water was chosen as a polar, and di-iodomethane as a non-polar liquid to compare the surface properties of the hybrid polymer layers containing different commercially available long chain perfluoroalkyl silanes (1 mol% each) with those of other surfaces. The measured contact angles and the surface energy data of the solids calculated according to the Owens and Wendt equation [24] are given in Table 6.3. The increase in the contact angles as a measure for a clearly reduced wettability can be observed with both test substances.

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6.11 XPS depth profile of the chemical elements silicon, aluminium and fluorine of hybrid sol–gel layers containing different quantities of the long chain perfluoroalkyl silane C8F17C2H4Si(OR)3: (a) 5 mol-%; (b) 1 mol-%

The coating system which was best suited in terms of anti-adhesive properties was to be tested for its colourability with dyes and pigments to provide for decorative effects, where required. Since the lowest value for the clear coating was found in compositions containing C8F17C2H4Si(OR)3, these systems were selected to be imparted with different colourants. The contact angles of the coloured nanocomposite layers towards water and di-iodomethane as well as the surface energy values calculated from them are also included in Table 6.3. Soluble dyes did not affect

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Table 6.3 Contact angle measurements (static values) on top of different surfaces as well as on clear and tinted hybrid sol–gel layers containing commercial long chain perfluoroalkyl silanes; surface energy values were calculated in accordance with Owens and Wendt [24] Surface or hybrid sol–gel layer (1 mol-% CnF2n+1C2H4Si(OR)3)

Colourant

PTFE Stainless steel Glass Hybrid sol–gel layer (without perfluoroalkyl silane) n=6 n = 6–10 n=8 n=8 n=8 n=8

– – – –

98 57 20 73

75 50 45 46

20.3 45.0 68.5 38.5

– – – Soluble dyes Organic pigments Inorganic pigments

96 102 103 106 105 98

76 90 90 91 83 83

20.3 13.8 13.5 12.7 16.0 17.0

Contact angle Contact angle (degrees) (degrees) H2O CH2J2

Surface energy (mN/m)

the surface energy. Incorporation of pigments, however, resulted in slightly increased surface energies, mainly due to a small reduction in the oleophobic nature (inorganic pigments slightly diminish the hydrophobic character). Perfluoroalkyl silanemodified clear and coloured layers showed similar properties compared to the basic composition specified in Table 6.2. The anti-adhesive effect becomes obvious when trying to write on a nanocomposite layer modified by a long chain perfluoroalkyl silane with a commercially available non-permanent overhead marker (Fig. 6.12). The surface reduces the wettability

6.12 Wettability of partially coated stainless steel substrates by a (non-permanent) overhead marker (left: transparent layer; right: coloured layer; both layers are equipped with a long chain perfluoroalkyl silane; C8F17C2H4Si(OR)3)

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towards the ink to such an extent that no homogenous film can be formed, and the ink contracts into droplets. The same effect is apparent on top of the coloured sol–gel layers. In contrast, conventional metal as well as glass or plastic surfaces are wetted homogeneously by this ink. Stainless steel surfaces coated with nanocomposite layers containing perfluoroalkyl silanes also have a clear improvement in their abrasion resistance (Fig. 6.13; variation of haze <0.1% after 100 cycles in the taber abraser test according to DIN 52347, determined in reflection). In contrast, uncoated stainless steel surfaces are much more sensitive towards scratching (increase in haze >5% after 100 cycles in the taber abraser test). Hence, a twofold protective performance is given for highly glossy metal surfaces by the hybrid polymeric nanocompounds. 6.5

Conclusions and outlook

As demonstrated, significant progress has been made in developing hybrid sol–gel materials via chemical nanotechnology, and in their application on metals. These materials display good durable adhesion when applied to various metal substrates and, because of their inorganic backbone, they exhibit high hardness and abrasion resistance. The coating materials can be adjusted to different requirements by varying their chemical composition and/or the process parameters. In addition, various functions may be combined in a single material, e.g. corrosion protection and mechanical protection or anti-adhesive properties and colourability and thus produce multifunctional coatings. Furthermore, potential substitutes for hexavalent chromium-based surface treatments for aluminium alloys are accessible on the basis of these nanocomposite materials. This kind of pre-treatment provides the organic coatings with superior adhesion to the aluminium metal surface and is responsible for the good corrosion resistance

6.13 Taber abraser test according to DIN 52347 (100 cycles) on top of partially coated stainless steel substrates

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of the coated aluminium. At present, hybrid polymers are being investigated as non-chromate pre-treatments for galvanised steel surfaces. And moreover, there will be trials to embed corrosion inhibitors (e.g. cerium compounds) into the hybrid polymer matrix to achieve ‘self-healing’ effects. Acknowledgements The investigations concerning chromate-free pretreatment methods have been funded by the Federal Ministry of Economics, procured by AiF (Arbeitsgemeinschaft industrieller Forschungsvereinigungen) and coordinated by DFO (Deutsche Forschungsgesellschaft für Oberflächenbehandlung e.V.). References 1. K.-P. Müller, Lehrbuch Oberflächentechnik, Vieweg Verlag, Braunschweig/Wiesbaden, 1996. 2. K.-H. Baumann, Korrosionsschutz für Metalle, Deutscher Verlag für Grundstoffindustrie, Leipzig, 1993. 3. K. Roths, J. Gochermann, K. Albinsky, K. Bonhoff, C. Borchers, R. Frenzel, S. Kölbl, U. Mock, D. Ondratschek, A. Patyk, A. Pflug (Autor), P. Plagemann (Autor), A. Schorb (Autor), B. Schmidt, S. Speith, V. Stenzel, B. Szyszka, P. Uhlmann, B. Voit, M. Weissenberger-Eibl, B. Wendler, Ch. Wieczoreck and D. Zickermann, Forschungsagenda Oberfläche. Analyse des Innovations- und Nachhaltigkeitspotenzials im Bereich der Oberflächenbehandlung, DFO Service, Neuss, 2006. 4. http://fao.dfo.info/vorstellung/Projektkurzbeschreibung_Forschungsagenda_berflaeche. pdf: Last accessed 1 June 2007. 5. P. Uhlmann, R. Frenzel, B. Voit, U. Mock, B. Szyszka, B. Schmidt, D. Ondratschek, J. Gochermann and K. Roths, Prog. Org. Coat., 58 (2007), 122. 6. J. Kron, K. Rose, K.-J. Deichmann, G. Schottner and K.-H. Haas, GAK, 56 (2003), 110. 7. J. Kron, K.-J. Deichmann, K. Rose, G. Schottner and K.-H. Haas, in Proceedings of the 1st Korrosionsschutz-Symposium, Korrosionsschutz durch Beschichtungen in Theorie und Praxis. Technische Akademie Wuppertal e. V., St. Goar, 2005. 8. EP01197539A2, EP 1310535, US06541563, US10292272. 9. Sol-Gel Processing of Ceramics and Glass, 2006; http://www.bccresearch.com/avm/ AVM016E.asp: Last accessed 1 June 2007. 10. S. Sakka, Handbook of Sol-gel Science and Technology: Processing, Characterization, and Applications. Springer, Boston, 2005. 11. http://www.atsdr.cdc.gov/toxprofiles/tp7.html. Last accessed 2 June 2007. 12. Directive 2000/53/EC of the European Parliament and of the Council of 18 September 2000 on end-of life vehicles; OJ L 269, 21.10.2000, 34–43; incl. Council Decisions 2002/525/ EC; 2005/63/EC; 2005/437/EC; 2005/438/EC; 2005/673/EC. 13. Directive 2002/95/EC of the European Parliament and of the Council of 27 January 2003 on the restriction of the use of certain hazardous substances in electrical and electronic equipment; OJ L 37, 13.2.2003, 19–23; incl. Council Decisions 2005/618/EC; 2005/717/EC; 2005/747/EC; 2006/310/EC; 2006/690/EC; 2006/691/EC; 2006/692/EC. 14. M. Bastian, DFO-Proceedings, Nr. 46, Leichtmetallanwendungen, Düsseldorf, 2001, 26. 15. U. Jüptner, Proceedings, Technische Fachtagung ‘Die industrielle Anwendung der Sol-Gel-Technologie’. Technische Akademie Wuppertal e. V., Würzburg, 2006. 16. P. Schubach, JOT, 7 (2000), 46. 17. Y. L. Leung, Y. P. Yang, P. C. Wong, K. A. R. Mitchell and T. Foster, J. Mater. Sci. Lett., 12 (1993), 844. 18. A. Sabata, W. J. Van Ooji and R. J. Koch, J. Adhes. Sci. Technol., 7 (1993), 1153.

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19. A. Klukowska, Thesis, University of Würzburg, Wuerzburg, Germany, 2004. 20. J. Kron, G. Schottner and K.-J. Deichmann, Thin Solid Films 392 (2001), 236. 21. J. Kron, K.-J. Deichmann, K. Rose, G. Schottner and K.-H. Haas, Proceedings 7. Fachforum ‘Schichten auf Glas’. OTTI-Technik Kolleg, Regensburg, 2004, 240. 22. International Workshop on Glasses, Ceramics, Hybrids and Nanocomposites from Gels, since 1981. 23. Journal for Sol-Gel Science and Technology, Kluwer Academic Publishers, since 1993. 24. D. K. Owens and R. C. Wendt, J. Appl. Polym. Sci., 13 (1969), 1741.

7 Hybrid sol–gel/conducting polymer coatings: self-healing coatings for the corrosion protection of aerospace alloys R. Akid, M. Gobara and H. Wang Materials and Engineering Research Institute, Sheffield Hallam University, Howard Street, Sheffield S1 1WB, UK [email protected]

7.1

Introduction

Corrosion causes loss of valuable products, affects operational safety, product reliability and may cause plant shutdown. It has been estimated that corrosion costs the UK economy between 3% and 4% of Gross National Product [1]. After ferrous alloys, aluminium is the most widely used engineering metal. Aluminium alloys are used in cars, aircraft, construction, engineering structures and many other industrial applications. Pure aluminium metal does not have suitable strength for most applications and therefore alloying elements, such as copper and magnesium, are added to improve the mechanical properties of the alloy. The 2xxx series alloys (Al/Cu) are very important in the aerospace industry; however, due to the presence of anodic and cathodic intermetallic particles, these alloys are not sufficiently resistant to localised corrosion to be used for external application without protection when exposed to humid or aggressive media such as marine environments. Chromate(VI)-based pre-treatments have historically been used to protect aluminium alloys. When the Cr(VI) species comes into contact with an aluminium substrate, a complex chromium oxide film is formed which provides excellent corrosion protection. Unfortunately, due to their toxicity and adverse environmental impact, these types of pre-treatment will be banned within Europe and North America. Many reports have shown that chromium(VI) is a carcinogen, and can cause kidney and liver damage, and even death [2,3]. Hence, there is an urgent need to provide corrosion protection systems for aluminium, and other active systems, that are accepted as being ‘environmentally-compliant’. A major advantage of hybrid sol–gel systems over conventional sol–gel coatings is that crack-free films with controlled composition and morphology can be formed from organic and inorganic oxide precursors at lower temperatures with respect to other ceramic-based systems. It is well recognised that coatings contain micropores, cracks, and areas of low cross-link density which provide a path for diffusion of corrosive species such as water, oxygen and chloride ions to the coating/metal interface [4–6]. Given this problem, a second advantage in applying a silica sol–gel coating to an aluminium alloy is that it can offer corrosion protection through the formation of a Si–Al rich intermediate reaction layer, which acts as an ‘inert’ barrier for the diffusion of oxygen to the metal surface thereby impeding and delaying the onset of corrosion [7]. 119

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Polyaniline (PANI) is a conductive polymer which has unique electrical and optical properties, moreover, it is relatively cheap, easy to synthesise and very stable under a wide variety of experimental conditions [8]. These polymers have some properties that may provide additional functionality other than corrosion protection, notably adsorption of electromagnetic waves. Since DeBerry [9] reported that polyaniline provided corrosion protection to stainless steel, there have been a large number of investigations using polyaniline as an additive in corrosion protection systems. Whilst the majority of research interest in using polyaniline for corrosion protection has been in combination with ferrous alloy substrates, there are a few reports concerned with the use of polyaniline for the corrosion protection of aluminium [10–13]. In the present study, polyaniline is combined with a silica sol–gel to offer a coating having the corrosion resistance benefits of polyaniline with the mechanical properties of the silica sol–gel. 7.2

Experimental

Chemical oxidation of aniline was performed by oxidation with ammonium persulphate in 1 M HCl solution; 5.0 ml of 0.107 M aniline was dissolved in 300 ml of 1 M HCl at 0–5ºC and stirred for 1 h. A solution of 5.6 g ammonium persulphate in 100 ml of 1 M HCl was added dropwise to the aniline solution over a period of 15 min with vigorous stirring. The solution was then continually stirred at 0–5ºC for 6 h. The precipitate was collected with a Buchner funnel and washed with four portions of 50 ml 1 M HCl. The precipitate was dried at 60ºC for 24 h, and the HCl doped PANI (Emeraldine Salt, ES) was obtained as a green powder. The basic form of PANI (Emeraldine Base, EB) was obtained by stirring the ES with 1 M NH4OH for 3 h followed by filtration and drying. The PANI sample was dissolved in Nmethylpyrrolidinone (NMP) solvent to form a 3.5% solution. Aluminium alloy 2024-T3 samples (obtained as Q-panels) were cleaned with deionised water followed by acetone and then left to dry for 30 min at 60ºC before applying the coating. The basic form of PANI (EB) was dissolved in NMP to give a 3.5% PANI solution. An organic-inorganic hybrid silica-based sol, catalysed by nitric acid, was prepared and used in this study; see Ref. 7 for further details. The PANI solution was added to the sol–gel liquid and stirred to form the PANI doped-sol. The PANI-doped sol was prepared in different relative volume concentrations as follows (PANI/sol): 1:1, 1:4, 1:8, 1:10, and 1:16. The coating was then applied to one side of the metal surface using a spray method. The coated metal was oven dried in air at 70ºC for 16 h. Corrosion tests were carried out in a three-electrode type cell using the sample (working electrode), a ‘saturated’ calomel reference electrode and a platinum counter electrode. Electrochemical impedance measurements were obtained at the measured Eocp values applying ±10 mV perturbation, in the frequency range from 3×104 Hz to 10−2 Hz. Electrochemical corrosion measurements were performed in 3.5 wt% NaCl (pH 6.8) solution at room temperature and in open air with an ACM instrument with the 3.5 wt% NaCl solution being renewed every 10 days. Additional corrosion testing using the above procedure was conducted in 3.5 wt% NaCl with acidic (HCl) and alkaline (NaOH) buffer solutions at pH 3.5 and 9.2, respectively. For the scratch test, electrochemical impedance spectroscopy (EIS) data were recorded for samples immersed in the corrosive solution for 5 days. Samples were

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then damaged using a scratch technique. The scratch was prepared using a razor blade giving a scratch width and length of about 100 μm and 5 cm, respectively which penetrated down to the substrate. After scratching, the sample was immersed in a fresh solution and EIS measurements carried out periodically. Salt spray testing was carried out at 35ºC within a 5% NaCl spray at 100% humidity for 500 h. Samples were checked every 48 h and images were recorded to compare the performance of samples throughout the test period. To assess adhesion, a Sellotape test was carried out as follows: the samples were scratched using a razor blade to form a cross-hatch pattern. Sellotape tape was applied to the cross-hatch and fixed well before peeling it from the surface. The Sellotape was examined for signs of detachment of the coating. 7.3 7.3.1

Results and discussion Short-term immersion testing

The electrochemical response of the sol–gel coated AA2024 was evaluated using EIS within a 3.5% NaCl solution for a typical coating thickness of 15 μm, as shown in Fig. 7.1. The coated AA shows high impedance (2×106 ohm.cm2) after a day of immersion; however, the impedance decreases with increasing time of immersion. This decrease in resistance correlates with the appearance of pitting after 4 days of immersion (figure not shown here). Figure 7.2 shows the behaviour of the sol–gel coating after 16 days, where the number of pits has not changed significantly from their appearance after 4 days; this result indicates that the silica sol–gel coating alone has limited corrosion protection for Al 2024 alloy. The initial high impedance of the coated sample reflects the hydrophobic nature of the coating [14], as seen around 104 Hz in the EIS phase diagram (Fig. 7.3) which shows the hydrophobicity of the coating decreasing with increasing time of immersion. After 1 day of immersion, there is only one time constant at 104 Hz relating to the coating. After 5 days, a second time constant is observed at 102 Hz, which is

7.1 Impedance of the sol–gel coated sample in 3.5% NaCl solution

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7.2 Image of pits on sol–gel coated sample after 16 days of immersion

7.3 Phase diagram of the sol–gel coated sample in 3.5% NaCl solution

representative of an aluminium oxide layer suggesting ingress of the electrolyte through to the metal substrate. Finally, after 16 days, a further time constant appears at 0.2 Hz which represents corrosion product formation and breakdown at the coating/metal interface. Although full details are not presented in this paper, coatings having PANI/sol–gel ratios of 1:16 and 1:12 were prepared and evaluated. These coating ratios showed pitting after 5 and 7 days, respectively, however, the degree of pitting was less than

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that found for the sol–gel only coated sample. PANI/sol–gel coatings prepared with concentration ratios of 1:1, 1:4 and 1:8 showed stable impedance values (as shown in Figs. 7.4–7.6) for up to 4 weeks of immersion in the corrosive solution with no signs of pitting. The impedance of 1:8 PANI/sol–gel within the first month was greater than that of the 1:4 and 1:1 compositions. This may in part be due to the thickness differences between the individual coatings, as shown in Fig. 7.7. Typical thickness values were: 12–13 μm, 9–10 μm and 5–6 μm for the 1:8, 1:4 and 1:1 ratios, respectively.

7.4 Impedance of the PANI/sol–gel 1:8 coated sample in 3.5% NaCl

7.5 Impedance of the PANI/sol–gel 1:4 coated sample in 3.5% NaCl

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7.6 Impedance of the PANI/sol–gel 1:1 coated sample in 3.5% NaCl

7.7 SEM of cross-sections of PANI/sol–gel: (a) 1:8, (b) 1:4 and (c) 1:1

7.3.2

Prolonged immersion testing

For prolonged immersion (up to 18 months), the 1:8 PANI/sol–gel ratio samples were chosen for testing as it was recognised that high sol–gel content improves the

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mechanical properties of the coating, a requirement where the coating may be subject to bending or stretching. Figure 7.8 shows the impedance behaviour during prolonged immersion. After 1 month of immersion, the hydrophobic nature of the coating decreases, as observed in a slight decrease in impedance within the frequency range of 10–1–10–2 Hz. However, the coating resistance remains stable above a value of 106 ohm.cm2 for a further 14 months. Physical examination of the coating showed no signs of pitting or delamination during this prolonged period, as shown in Fig. 7.9. 7.3.3

Influence of solution pH on coating performance

Individual samples of 1:8 ratio coated AA were immersed separately in 3.5% NaCl (pH 3.5) or (pH 9.2) solutions. EIS data were recorded to evaluate the corrosion

7.8 Impedance of the PANI/sol–gel 1:8 coated sample in 3.5% NaCl

7.9 Image of PANI/sol–gel 1:8 coated sample after 15 months in 3.5% NaCl solution (note black spots are particles of PANI not pitting)

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performance at each pH value as a function of time. Figure 7.10 shows the impedance curve of the coated sample in acidic solution. As seen from Fig. 7.10, the impedance of the coated sample is one order of magnitude higher than that of the bare sample. According to the Pourbaix diagram, pure aluminium is passive in the pH range of 4–8 [15], however, as shown from Fig. 7.10, the coated sample is stable (105 ohm.cm2), and one order of magnitude higher than that of the bare sample following 2 months of immersion indicating protection of the substrate over this time. The impedance curve of the 1:8 ratio sample in 3.5% NaCl alkaline solution (pH 9.2) is shown in Fig. 7.11. It is clear from this figure that the bare metal sample has a poor resistance (less than 104 ohm.cm2 after 24 h of immersion) in this corrosive

7.10 Impedance of bare and PANI/sol–gel coated AA2024 in 3.5% NaCl (pH 3.5) solution

7.11 Impedance of bare and PANI/sol–gel coated AA2024 in 3.5% NaCl (pH 9.2) solution

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alkaline solution. Conversely, the PANI/sol–gel coated sample remains stable over a 45-day period at 105 ohm.cm2, without evidence of pitting or delamination of the coating. The EIS results for both acid and alkaline solutions suggest that PANI/ sol–gel coating can offer corrosion protection within the pH range of 3.5–9.2. 7.3.4

Scratch test (self-healing assessment)

A 1:8 ratio sample was immersed for 5 days in 3.5%NaCl solution before scratching and immersion in a fresh solution of the same composition. EIS data representing the coating behaviour before and after scratching is shown in Fig. 7.12a,b. It is clear from the impedance curves that when the coated metal is scratched, the corrosion resistance drops by around one order of magnitude. However, with prolonged

7.12 (a) Impedance of 1:8 PANI/sol–gel coated AA2024 in 3.5% NaCl solution before and after scratching. (b) Phase angle of 1:8 PANI/sol–gel coated AA2024 in 3.5% NaCl solution before and after scratching

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immersion, the impedance value returns towards the pre-scratched value regaining about half an order of magnitude resistance after 45 days. This result indicates that PANI/sol–gel coating exhibits ‘self-healing’ type behaviour. The phase plot (Fig. 7.12b) shows that, after being scratched, the sample has two time constants, at 2×102 and 1×10–1 Hz, which may be related to the coating and metal/coating interface, respectively. On increasing the immersion time, the high frequency time constant moved towards lower frequency and its peak decreased with immersion time. The low frequency time constant peak increased with immersion time indicating an increase in capacitance of the interface. This was combined with an increase in impedance in the same frequency region. Moreover, the capacitive behaviour was clearly defined with immersion time; such behaviour is not seen with inert barrier coatings. This may be due to the PANI ‘self-healing’ property, which improved with immersion time. On physical examination, the PANI/sol–gel coated samples showed neither undercutting nor blistering. Furthermore, when the Sellotape test was applied, no loss of adhesion was noted, rather accumulation of white aluminium oxide (arrows) was observed at the scratch line as shown in Fig. 7.13. An understanding of the mechanism by which the self healing ‘repair’ process takes place was investigated using Fourier transform infra-red (FTIR) analysis. Figure 7.14 shows the molecular configurations of the reduced and oxidised forms of polyaniline and the corresponding FTIR wavenumbers. FTIR analyses of samples of polyaniline on aluminium before and after immersion in 3.5% NaCl, were compared with equivalent samples immersed in solution. Comparison was also made with samples immersed and then air dried for 1 week. The results of this analysis are given in Fig. 7.15. Here the relationship (ratio) of the reduced (A) to oxidised (B) forms was measured for the four different samples. This ratio reflects the oxidised state of the surface and can be correlated with the corrosion behaviour of the system.

7.13 Scratched PANI/sol–gel coated sample after 5 months in 3.5% NaCl

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7.14 Molecular structures of different forms of polyaniline and the respective FTIR wave number for the reduced (A) and oxidised (B) forms

7.15 FTIR spectrum of PANI coated AA2024 before and after immersion in 3.5% NaCl solution

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According to the results in Fig. 7.15, when the EB form of PANI was immersed in NaCl solution, the relative concentration of the reduced form (A) to the oxidised form (B) increased. This can take place in one of the following two ways: a) The oxidised form (B) is converted to the reduced form (A) by the following reaction: PANI (B) + e– ⇒ PANI (A)

[7.1]

In this case, the supply of electrons is provided by the oxidation of Al to aluminium ion. Al ⇒ Al3+ + 3e–

[7.2]

In this case, Al will continue to corrode but at a higher corrosion rate than without the PANI coating which did not occur. Alternatively, b) The oxidised form (B) undergoes a reaction, for example: 1) PANI undergoes cross-linking as shown in Fig. 7.16 According to Lee et al. [16], PANI forms cross-linked molecules above 80°C, moreover, the cross-linked molecule cannot be re-oxidised. From Fig. 7.15, it can be seen that the ratio of A/B decreased after 1 week following drying in air. This result indicates that the mechanism of cross-linking proposed above does not take place. An alternative mechanism is suggested:

7.16 Proposed cross-linking scheme of PANI (EB) [16]

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2) Reaction of the oxidised form (B) with the Al substrate to form an Al-PANI complex. A theoretical study by Calderone et al. [17] concluded that the oxidised form of PANI is more reactive with metal atoms than the reduced form of PANI. In this case, any Al-PANI complex formed will act as a protective layer causing an increase in impedance of the coating. Moreover, the reduced form of PANI is oxidised in air at room temperature and this reaction can take place many times [18]. 7.3.5

Salt spray performance

Salt spray testing (SST) was performed using unscratched bare and scratched sol–gel only and 1:8 PANI/sol–gel coated samples. Corrosion of the bare Al alloy surface takes place within 48 h, as shown in Fig. 7.17a. A performance improvement was noted for the sol–gel only coating where the number of pits observed on the surface

7.17 Images of SST samples: (a) Bare sample after 48 h; (b) sol–gel coated sample after 72 h and (c) 1:8 PANI/sol–gel after 350 h

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was less than those confined to the scratch site (Fig. 7.17b). The greatest protection was observed with the PANI/sol–gel 1:8 ratio coated samples which did not show any pitting on the surface or at the scratch site, for up to 500 h of salt spray testing, as shown in Fig. 7.17c. Adhesion of the coating was also assessed following SST using a simple ‘Sellotape’ test. No undercutting or blistering of the 1:8 ratio coating was observed. 7.4

Conclusions

A series of PANI/sol–gel coating formulations have been applied to an AA2024-T3 substrate and evaluated using electrochemical and salt spray tests. Comparison of the corrosion protection performance of these coatings was made with bare and sol–gel only coated samples. The following conclusions can be drawn: Silica sol–gel can offer limited corrosion protection to AA2024-T3 when subjected to SST and immersion testing. Immersion in neutral 3.5% NaCl solution showed that a 1:8 ratio PANI/sol–gel coating could protect AA2024-T3 for 18 months of immersion without pitting or delamination of the coating. Moreover, this coating can protect AA2024-T3 in both acidic and alkaline media. Scratch testing showed that a PANI/sol–gel coating could offer self-healing properties to the AA2024-T3. Self healing was ascribed to transformation of the oxidised and reduced form of PANI and the possible formation of an Al-PANI complex. Acknowledgements The authors acknowledge the Materials and Engineering Research Institute at Sheffield Hallam University for the provision of facilities and to the Egyptian Government for providing funding for M. Gobara. References 1. Cost of Corrosion. The Institute of Materials, Minerals and Mining, http://www.iom3. org/divisions/surface/corrosion/index.htm. Last accessed: 2. M. Costa and C. B. Klein, Crit. Rev. Toxicol., 36(2) (2006), 155–163. 3. L. Assem and H. Zhu, Chromium Toxicological Overview. Protection Agency, UK. www. hpa.org.uk/web/HPAwebFile/HPAweb_C/1194947362170 (Health, 2007), p. 14. 4. X. F. Yang, D. E. Tallman, V. J. Gelling, G. P. Bierwagen, L. S. Kasten and J. Berg, Surf. Coat. Technol., 140(1) (2001), 44–50. 5. M. L. Zheludkevich, I. Miranda Salvado and M. G. S. Ferreira, Mater. Chem., 15 (2005), 5099–5111. 6. A. S. Hamdy and D. P. Butt, Surf. Coat. Technol., 201(1–2) (2006), 401–407. 7. H. Wang and R. Akid, Corros. Sci., 49(12) (2007), 4491–4503. 8. E. M. Geniès, A. Boyle and M. Lapkowski, Synth. Met., 36(2) (1990), 139–182. 9. D. W. DeBerry, J. Electrochem. Soc., 132(5) (1985), 1022–1026. 10. R. Racicot, R. Brown and S. C. Yang, Synth. Met., 85(1–3) (1997), 1263–1264. 11. D. E. Tallman, Y. Pae and G. P. Bierwagen, Corrosion, 56(4) (2000), 401–410. 12. A. J. Epstein, J. A. O. Smallfield, H. Guan and M. Fahlman, Synth. Met., 102(1–3) (1999), 1374–1376.

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13. L. Cecchetto, R. Ambat, A. J. Davenport, D. Delabouglise, J. Petit and O. Neel, Corros. Sci., 49(2) (2007), 818–829. 14. W. A. Daoud, J. H. Xin and X. M. Tao, J. Am. Ceram. Soc., 87(9) (2004), 1782–1784. 15. C. Vargel, M. Jacques, and D. M. P. Schmidt, The Corrosion of Aluminium. Elsevier, Amsterdam, 2004, 81. 16. Y. M. Lee, J. H. Kim, J. S. Kang and S. Ha, Macromolecules, 33(20) (2000), 7431–7439. 17. A. Calderone, R. Lazzaroni and J. L. Brédas, Phys. Rev. B, 49(20) (1994), 14418–14426. 18. S. Yang, R. Brown and J. Sinko, Eur. Coat. J., 11 (2005), 48–54.

8 Corrosion performance of nanoparticle-containing polyaniline films on AA3105 aluminium alloy O. Zubillaga and F. J. Cano INASMET, Mikeletegi Pasealekua 2, San Sebastian, Spain [email protected]

A. M. Cabral and P. J. Morais ISQ-Instituto de Soldadura e Qualidade, Tagus Park, 2740-120 Porto Salvo, Portugal

I. S. Molchan, P. Skeldon and G. E. Thompson Corrosion and Protection Centre, School of Materials, The University of Manchester, PO Box 88, Manchester, M60 1QD, UK

T. Schmidt and M. Schem Leibniz Institute of Neuen Materialien, Im Stadtwald, D-66123 Saarbrücken, Germany

8.1

Introduction

The widespread use of aluminium alloys in applications ranging from transport to architecture demands surface treatments for their corrosion protection [1,2]. Anodising of aluminium is a widely studied and accepted method that consists of electrochemically forming a protective aluminium oxide layer [2]. The growth, morphology, composition and performance of anodic films have been considered in the literature, for both pure aluminium and aluminium alloys [2–7]. Furthermore, sealing mechanisms of the porous aluminium oxide layers have also been reported [6–9]. Various alternatives have been proposed in recent years as replacements for anodising in chromic acid for protection of aluminium alloys from corrosion. For example, Moutalier and colleagues [10–12] studied the incorporation of cerium(IV) and molybdate ions as inhibitors in the anodic films formed in sulphuric acid media. An improvement in corrosion protection was reported for AA2024-T3 aluminium alloy compared with anodising in the absence of such ions. Further, Kamada et al. [13] reported the incorporation of oxide nanoparticles into alumina films, with anodic oxidation of the aluminium and electrophoretic deposition of nanoparticles proceeding in a single step. Alumina films containing SiO2 nanoparticles showed an increase in capacitance and improved corrosion protection of the substrate. The incorporation of SiO2 nanoparticles into the nanosized pores of anodic alumina films was also examined with the surface charge and size of the nanoparticles considered as being significant in enabling nanoparticles to migrate to the substrate during anodising with subsequent incorporation into the anodic film [14]. 134

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A further approach to the provision of chromate-free corrosion protection solutions is the use of conducting polymers that may be electropolymerised on aluminium substrates by anodic oxidation. However, porous anodic alumina films may grow in the acidic solutions commonly used for the electrosynthesis of conducting polymers, thereby acting as a barrier for electron transfer and, thus, inhibiting the polymerisation process. However, successful electrosynthesis of polyaniline and polypyrrole on aluminium and its alloys has been reported. Beck and Husler [15] showed that the pores of the anodic alumina layer play an important role as nucleation centres in the electropolymerisation of pyrrole on aluminium. Huerta-Vilca et al. [16] also concluded that primary polymerisation of aniline on aluminium surfaces proceeds in the pores of the anodised surface. Naoi et al. [17] proposed a mechanism for the simultaneous formation of a barrier alumina film and polypyrrole over the alumina surface, where the initially formed alumina film contained cracks in which the hydrophilic groups of the surfactant electrolyte are attached; the hydrophobic groups form a micelle that hosts the polypyrrole molecules. The electropolymerisation of pyrrole in the micelles forms electronically conducting paths of polypyrrole that extend from the aluminium electrode to the surface of the alumina film. Such conducting paths allow current flow and formation of polypyrrole layers on the surface of the alumina film. Tsai et al. [18] proposed a three-stage mechanism, similar to that of Naoi et al.; the first stage consisted of the formation of porous anodic alumina and the nucleation of polypyrrole within the pores; this was followed by the propagation of the polymer over the alumina layer and, finally, over-oxidation of the polypyrrole, located within the pores, proceeded. Incorporation of inorganic fillers into conducting polymer matrices has also been studied, utilising both chemical [19–23] and electrochemical [19,24–27] polymerisation processes. For example, TiO2 [22,24–27], ZrO2 [21,23,26] and SiO2 [26] particles have been incorporated in polyaniline and polypyrrole matrices. Zhu and Iroh [26] electrosynthesised polyaniline in the presence of such nanoparticles on an AA2024 aluminium alloy. Potentiodynamic polarisation tests suggested a positive effect of the fillers on the corrosion resistance provided by polyaniline coatings, although no information on the mechanism was provided. Lenz et al. [27] electrosynthesised polypyrrole on mild steel in the presence of TiO2 particles; the corrosion resistance was examined by measuring the release of iron to the 3.5% NaCl solution and the associated weight loss of the specimen. The polypyrrole/TiO2 composite films showed improved performance compared with the polypyrrole films; this was attributed to the low porosity of polypyrrole/TiO2 films due to the filling of the polymer pores by TiO2 particles. Given the ability of incorporating nanoparticles into porous anodic films and conducting polymer matrices, with the latter forming within the pores, the incorporation of both TiO2 nanoparticles and polyaniline in single step anodising is considered here. The corrosion performance of the coated aluminium alloy has been examined by salt spray testing and by electrochemical impedance spectroscopy. 8.2 8.2.1

Experimental Materials and treatment

An AA3105 aluminium alloy of 0.5 mm thickness was used as the substrate; the alloy composition is given in Table 8.1. Specimens of dimensions 50×100 mm were cut from the as-received sheets.

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Table 8.1 Chemical composition of the AA3105 aluminium alloy provided by the supplier Element (wt.%)

Si

Fe

Cu

0.298 0.629 0.044

Mn

Mg

Zn

0.437

0.350

0.025

Ti

Cr

Ni

Al

0.018 0.022 0.013 98.160

Before anodising, the specimens were degreased with acetone, followed by alkaline etching in Turco 4215NCLT for 10 min at 50ºC and, finally, desmutting in Turco Smut Go for 5 min at room temperature. Anodising was undertaken at a constant voltage of 8 V in 0.5 M oxalic acid with 0.1 M aniline and 1 g/l of dispersed TiO2 nanoparticles of 5–10 nm size (measured by HR-TEM); the cell was maintained under a nitrogen atmosphere with continuous stirring of the pH 1 electrolyte. Specimens were also anodised in a similar electrolyte without nanoparticles, for comparison. 8.2.2

Coating morphology and composition characterisation

The surface morphology and composition of the coatings were examined by scanning electron microscopy with energy dispersive spectroscopy (SEM/EDS), using a JEOL JSM 5910LV microscope fitted with an INCA 300 spectrometer. The cross-sections of the coated substrates were also examined by SEM and transmission electron microscopy (TEM). For TEM examination, electron transparent sections were prepared using a Leica Ultracut UCT ultramicrotome. Final sections, of nominal thickness 15 nm, were generated with a diamond knife (Microstar) from the coated alloy that had been encapsulated previously in cured resin, and examined in a JEOL 200FX II transmission electron microscope. The coating composition was analysed by X-ray photoelectron spectroscopy (XPS) using a PHI 5500 Multitechnique System. Analysis was undertaken in an area of 0.8 mm diameter. The depth profiles were generated by sputtering the surface with an argon ion source of 4 keV energy. The depth values were calculated using a sputtering rate of 7 nm/min, obtained from a Si3N4 calibration standard. The error of the XPS measurement is 1% in the quantification of the content of each element. 8.2.3

Salt spray testing

Salt spray tests (SST) were carried out under standard conditions (EN ISO 9227 NSS) in a C&W/SF-1000-CCT chamber; the coated specimens were exposed for periods up to 1000 h. The progress of corrosion was evaluated from macrographs of the exposed surfaces that were recorded at different exposure times. 8.2.4

Impedance spectroscopy measurements

Electrochemical impedance spectroscopy (EIS) analysis was performed in 0.5 M NaCl solution at room temperature, using a Gamry REF 600 potentiostat. A threeelectrode cell was employed, with the variously treated AA3105 aluminium alloys as the working electrode, a saturated calomel reference electrode and a platinum wire counter electrode. Before EIS analysis, the open-circuit potential of the specimen was continuously monitored from immersion in the electrolyte for 30 min. The EIS data were acquired over the frequency region of 1 MHz to 0.001 Hz at 7.13 points per

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decade, with an ac voltage amplitude of 10 mV. The cell was formed by attaching a Perspex cylinder to the specimen surface, with an exposed area of 3.8 cm2. 8.3 8.3.1

Results and discussion Coating morphology and composition

Figure 8.1 shows scanning electron micrographs of the surface of the etched AA3105 alloy after electrochemical treatment in the acid solution containing aniline in the absence (Fig. 8.1a) and presence (Fig. 8.1b) of nanoparticles. Numerous cavities are evident in the micrographs that are associated with loss of second phase particles

8.1 Surface SEM micrograph of 3105 aluminium alloy anodised in oxalic media (a) with polyaniline and (b) polyaniline and TiO2 nanoparticles. In (b), the agglomeration of TiO2 nanoparticles is indicated with number 1 and some examples are marked with circles

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138

during the alkaline etching process. The cavities are randomly shaped, with dimensions from 1 to 8 μm, and population density of approximately 3.5x105 cm–2. Bright features appear on the surface of the anodic film formed in the presence of nanoparticles (Fig. 8.1b). They represent regions of increased nanoparticle content, indicating that particles form aggregates over the surface of the anodic film. The increased titanium content in the bright areas was confirmed by EDS. Cross-sections of the film generated on the alloy reveal a film thickness of about 2.7 μm (Fig. 8.2). The

8.2 Cross-section (a) TEM and (b) (SEM) micrographs. In (a): 1, alloy; 2, anodic film, 3, nanoparticle-rich region. In (b): 1, second-phase particles; 2, cavities of the anodic coating covered by a thin anodic layer; 3, thickness of the anodic layer

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transmission electron micrograph of the ultramicrotomed section of the treated alloy reveals the porous anodic film attached to the alloy substrate. A thin dark, particleenriched layer is present at the surface of the anodic film (Fig. 8.2a). In the scanning electron micrograph, locally disrupted areas that are associated with the presence of second phase Fe/Mn/Si intermetallics in the alloy are evident (Fig. 8.2b). In the zone with evidence of defects, at the substrate/anodic film interface, where intermetallics were removed during alkaline etching, a thinner anodic film is observed compared with the anodic film over the macroscopic alloy surface. Figure 8.3 shows transmission electron micrographs of selected regions of the anodic film located near the surface (Fig. 8.3a), within the film (Fig. 8.3b), and adjacent to the alloy/film interface (Fig. 8.3c). The anodic film reveals a porous morphology, with pores confined within alumina cells; the barrier layer is evident beneath the pore base together with the scalloped alloy/film interface. The cell and pore diameters are about 30 and 10 nm, respectively, with a barrier layer thickness of 11 nm. Titania nanoparticles are distributed in a thin, outer layer of several tens of nanometres thickness; the particles have diameters up to 10 nm. No particles are evident in the porous skeleton of the anodic film. An XPS depth profile (Fig. 8.4) of the particle-rich region indicates that an elemental percentage of titanium of about 11% at the outer surface, decreases gradually over a thickness of approximately 100–150 nm. The aluminium signal corresponds to the aluminium present in the aluminium oxide of the anodic film and reveals some defects in the polyaniline-TiO2 layer of the outer part of the anodic film. This aluminium is not related to the substrate aluminium, and thus it does not indicate the presence of defects that allow the electrons coming from the substrate. 8.3.2

Salt spray testing

The influence of TiO2 nanoparticles on the performance of the AA3105 aluminium alloy was assessed by salt spray testing (SST), as shown in Fig. 8.5. The alloy supporting the coating comprising the alumina film and polyaniline revealed corrosion products and pits over the surface for exposure times below 168 h and the surface coverage by corrosion products increased with time. Conversely, the coating with incorporated TiO2 nanoparticles did not reveal evidence of corrosion for exposure times of 1000 h. Furthermore, no corrosion was detected in the scribed area of the alloy that originally supported the coating with nanoparticles. 8.3.3

Electrochemical impedance spectroscopy

The EIS spectra of the AA3105 aluminium alloy supporting the anodic film with polyaniline in the presence and absence of TiO2 nanoparticles during immersion in 0.5 M NaCl solution for 30 min, 4 days and 8 days are displayed in Fig. 8.6. Spectra for the bare alloy are also shown for comparison. Spectra for the early times of exposure (Fig. 8.6b) show that the untreated alloy presents a time constant at frequencies around 1 kHz, that can be associated with the natural oxide layer present on the bare alloy. The alloy treated with polyaniline, shows a broad time constant, suggesting two overlapping time constants (at frequencies around 500 Hz and 5 kHz) that can be associated with the presence of the anodic porous aluminium oxide film and the polyaniline layer, respectively. Further, the alloy treated with polyaniline showed an impedance modulus one order of magnitude

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8.3 Cross-section TEM micrograph of 3105 aluminium alloy anodised in oxalic media with polyaniline and TiO2 nanoparticles. Selected regions of the anodic film located (a) near surface, (b) within the film, (c) and adjacent to the alloy/film interface

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8.4 XPS depth profile for the polyaniline-TiO2 layer and part of the anodised layer

higher than the bare alloy (Fig. 8.6a). The untreated alloy and the alloy treated with polyaniline reveal different behaviour, the treated alloy having a capacitive behaviour. However, this behaviour was progressively blurred with immersion time, and shifts in the time constant (≈90º) for lower frequency regions were observed for 5 h and 4 days of immersion (Fig. 8.6d). This behaviour is also evident in the abrupt decrease observed at low frequency impedance, from around 1x105 ohm.cm2 for 30 min of immersion to a value around 5x104 ohm.cm2 for increased times (Fig. 8.6a,c). For immersion times longer than 4 days, a time constant at lower frequencies was detected (Fig. 8.6d,f). The region of the spectra in which such time constant appears is usually ascribed to ionic diffusion processes at the substrate surface. Such processes are more important with time of immersion as a result of corrosion that progresses at the metal surface, with the formation of corrosion products. For increased immersion times (Fig. 8.6e,f), the behaviour is very similar to that evident for the bare substrate, indicating that the diffusion process gradually becomes more important. The EIS results obtained for the alloy supporting the coating with nanoparticles showed contrasting behaviour to that described previously. As observed in Fig. 8.6a, the low frequency impedance for the TiO2 containing coating is more than one decade higher than the impedance of the coating without nanoparticles, suggesting improved corrosion protection performance. After 4 days and 8 days of immersion in the NaCl solution (Fig. 8.6c,e), the low frequency impedance still remains around 1 decade higher for the TiO2 containing coatings. In this case, two time constants are observed (Fig. 8.6b,d,f); the time constant at higher frequencies (around 100 kHz) is associated with the TiO2 particle-rich thin layer and the constant evident at lower frequencies (around 50 Hz) is associated with the porous anodic aluminium layer. The time constant attributed to the outer thin layer of the coating is detected at a frequency of around 100 Hz for the coating without TiO2 nanoparticles whereas, for the TiO2 containing coating, this time constant is detected at higher frequencies of around 100 kHz; this suggests that the presence of TiO2 nanoparticles leads to a coating with lower capacitance. The time constant associated with the porous anodic film was detected

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8.5 Anodised specimens after 1000 h of salt spray exposure: (a) anodised layer with polyaniline and without TiO2 nanoparticles and (b) anodised layer with polyaniline and TiO2 nanoparticles

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Corrosion performance of nanoparticle-containing polyaniline films

8.6 Impedance modulus and phase angle evolution with immersion time: (a) impedance modulus at 30 min immersion time; (b) phase angle at 30 min immersion time; (c) impedance modulus at 4 days immersion time; (d) phase angle at 4 days immersion time; (e) impedance modulus at 8 days immersion time; (f) phase angle at 8 days immersion time

144

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at lower frequency values (around 50 Hz) compared with the coating without nanoparticles (around 500 Hz); this indicates that the presence of the TiO2 nanoparticles in the anodising electrolyte also influences the formation of the porous film. Further, for the TiO2 containing coating, the high frequency time constant does not vanish with time, as observed in the case of the alloy supporting the coating with polyaniline alone. Additionally, despite an initial decrease, a shift to higher frequencies is evident, which can be related to modifications of the TiO2 particle-rich layer in contact with the electrolyte. On the other hand, for longer immersion times, a time constant associated with diffusion processes was detected in the low frequency region (Fig. 8.6f). More precise analysis is possible from estimation of the coating capacitance and coating resistance; established approximation methods were used [28]. The equation C=–1/2πfZimag was used for the capacitance estimation, where f is the frequency of the maximum in the phase angle value and Zimag is the imaginary part of the impedance at this frequency. The data used in this case correspond to the time constants present in the higher frequency region, associated with the outer thin coating layer. The equation R=Zreal was used for the resistance estimation, where Zreal is the real part of the impedance at the frequency with minimum phase angle. In this case, the data used for the resistance estimation correspond to the minimum phase angles present at low frequencies. The variation of the resistance and capacitance of the alloy supporting the anodic film and polyaniline alone with immersion time in the chloride solution is shown in Fig. 8.7. The resistance reveals a rapid decrease from the commencement of immersion that is mirrored by an increase in capacitance. With increasing immersion time, the resistance recovered, possibly due to the presence of corrosion products and the capacitance decreased, probably reflecting a reduced area of corrosion through stifling by corrosion products. The capacitance values (Fig. 8.7a) for the TiO2 containing coatings range between 1.6x10–7 F/cm2 (30 min immersion time) and 3.9x10–8 F/cm2 (8 days immersion time). These values are more than two orders of magnitude lower than those observed for the coating without nanoparticles. Additionally, the coating without nanoparticles shows about one order of magnitude lower resistance values. Both the higher coating resistance and the lower capacitance values are evidence of the improved barrier properties of the TiO2 containing coatings. Concerning the initial resistance increase and capacitance decrease observed in the coatings with TiO2 nanoparticles, these are considered to be associated with some modification of the TiO2 particle-rich layer occurring in contact with the electrolyte. At present, this phenomenon is not well understood, and will be examined in more detail in further work. 8.4

Conclusions

Anodic alumina films with polyaniline and TiO2 nanoparticle-enriched regions on their surface were electrochemically synthesised on an AA3105 aluminium alloy in oxalic acid electrolyte containing dissolved aniline and dispersed nanoparticles. The resultant coatings comprised a porous anodic alumina film of about 2.7 μm thickness, with an outer hybrid polyaniline/TiO2 layer of several tens of nanometres thickness. Both EIS analysis and SST results revealed that the coatings with incorporated nanoparticles provide improved corrosion protection to the AA3105 aluminium

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8.7 Evolution of coating (a) capacitance and (b) resistance with immersion time

alloy compared with those without nanoparticles. In SST, aluminium specimens coated with alumina–polyaniline–TiO2 film did not reveal corrosion products after 1000 h of exposure, whereas corrosion products were detected for exposure times below 168 h for the specimens supporting coatings without TiO2 nanoparticles. The EIS analysis showed that the coatings containing TiO2 nanoparticles present lower capacitance and higher resistance values. The improved protection provided by coatings containing TiO2 nanoparticles is attributed to the TiO2 particle-rich thin film layer formed on the outer part of the coating that acts as a blocking barrier layer for the anodic porous aluminium oxide film. The role of the polyaniline and inhibiting properties of TiO2 nanoparticles will be the subject of further study by the authors. Acknowledgements The authors acknowledge the European Commission for the funding to carry out this work under the project MULTIPROTECT (Contract Nº NMP3-CT-2005-011783).

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References 1. J. R. Davis, Corrosion of Aluminum and Aluminum Alloys. First edition. ASM International, Materials Park, OH, USA, 1999. 2. A. W. Brace, The Technology of Anodizing Aluminium. Third edition. Interall S.r.l., Italy, 2000. 3. V. P. Parkhutik, Corros. Sci., 26(4) (1986), 295–310. 4. H. Habazaki, K. Shimizu, P. Skeldon, G. E. Thompson, G. C. Wood and X. Zhou, Corros. Sci., 39(4) (1997), 731–737. 5. M. A. Páez, T. M. Foong, C. T. Ni, G. E. Thompson, K. Shimizu, H. Habazaki, P. Skeldon and G. C. Wood, Corros. Sci., 38(1) (1996), 59–72. 6. A. Bautista, J. A. González and V. López, Surf. Coat. Technol., 154 (2002), 49–54. 7. V. Moutarlier, M. P. Gigandet, J. Pagetti and B. Normand, Surf. Coat. Technol., 182 (2004), 117–123. 8. F. Snogan, C. Blanc, G. Mankowski and N. Pébére, Surf. Coat. Technol., 154 (2002), 94–103. 9. M. J. Bartolomé, V. López, E. Escudero, G. Caruana and J. A. González, Surf. Coat. Technol., 200 (2006), 4530–4537. 10. V. Moutalier, M. P. Gigandet, L. Ricq and J. Pagetti, Appl. Surf. Sci., 183 (2001), 1–9. 11. V. Moutalier, M. P. Gigandet, J. Pagetti and B. Normand, Surf. Coat. Technol., 161 (2002), 267–274. 12. V. Moutalier, M. P. Gigandet, J. Pagetti and L. Ricq, Surf. Coat. Technol., 173 (2003), 87–95. 13. K. Kamada, M. Tokutomi, N. Enomoto and J. Hojo, J. Mater. Chem., 15 (2005), 3388– 3394. 14. K. Kamada, H. Fukuda, K. Maehara, Y. Yoshida, M. Nakai, S. Hasuo and Y. Matsumoto, Electrochem. Solid State Lett., 7 (2004), B24–B28. 15. F. Beck and P. Husler, J. Electroanal. Chem., 280 (1990), 159–166. 16. D. Huerta-Vilca, S. Regina de Moraes and A. de Jesus Motheo, Synth. Met., 140 (2004), 23–27. 17. K. Naoi, M. Takeda, H. Kanno, M. Sakakura and A. Shimada, Electrochim. Acta, 45 (2000), 3413–3421. 18. M. L. Tsai, P. J. Chen and J. S. Do, J. Power Sources, 133 (2004), 302–311. 19. R. Gangopadhyay and A. De, Chem. Mater., 12 (2000), 608–622. 20. D. C. Schnitzler and A. J. G. Zarbin, J. Brazil. Chem. Soc., 15 (2004), 378–384. 21. S. Sinha Ray and M. Biswas, Synth. Met., 108 (2000), 231–236. 22. K. Gurunathan and D. C. Trivedi, Mater. Lett., 45 (2000), 262–268. 23. A. Bhattacharya, K. M. Ganguly, A. De and S. Sarkar, Mater. Res. Bull., 31(5) (1996), 527–530. 24. D. M. Lenz, C. A. Ferreira and M. Delamar, Synth. Met., 126 (2002), 179–182. 25. Y. Liu, J. Huang, C. Tsai, T. C. Chuang and C. Wang, Chem. Phys. Lett., 387 (2004), 155–159. 26. Y. Zhu and J. O. Iroh, J. Adv. Mater., 34(4) (2002), 16–21. 27. D. M. Lenz, M. Delamar and C. A. Ferreira, J. Electroanal. Chem., 540 (2003), 35–44. 28. R. C. Cottis and S. Turgoose, Electrochemical Impedance and Noise. NACE International, Houston, TX, USA, 1999.

9 Advances in the selection and use of rare-earth-based inhibitors for self-healing organic coatings S. J. Garcia, J. M. C. Mol and J. H. W. de Wit Delft University of Technology, Department of Materials Science and Engineering, Mekelweg 2, 2628CD Delft, The Netherlands [email protected]

T. H. Muster, A. E. Hughes, T. Miller and T. Markley CSIRO Materials Science & Technology, Private Bag 33, Clayton South MDC, VIC 3169, Australia

J. Mardel CSIRO Molecular and Health Technologies, Private Bag 10, Clayton, VIC 3169, Australia

H. Terryn Vrije Universiteit Brussel (VUB), Department of Materials and Chemistry – META, Pleinlaan 2, 1050 Brussels, Belgium

9.1

Introduction

The corrosion of metals results in significant costs worldwide associated with repair and replacement [1–4]. These costs are typically 2% to 5% of gross domestic product (GDP) [1,4]. Metal corrosion is predominantly an electrochemical process that needs four elements to take place: an anode, a cathode, an electrolyte for ionic mobility, and an electrical connection between the anode and the cathode. If one or several of these elements is hindered, for example, through the use of corrosion inhibitors, cathodic protection mechanisms, or through the application of passive and/or active barrier systems such as paint, the electrochemical process will be slowed or, ideally, stopped. In the fight to reduce corrosion rates, organic primers are widely used, since they are both efficient and relatively cheap, and can be used in many different applications. Organic coatings can protect the substrate by one or several of the following mechanisms: 1. 2. 3. 4.

provide strong adherence to avoid delamination provide a barrier to minimise mass-transport carry sacrificial pigments carry inhibitive species (i.e. corrosion inhibitors).

No organic coating provides a perfect barrier against the ingress of moisture and oxygen. Therefore, for primer coatings that have a high performance demand, it is 148

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important that the inhibitive and sacrificial pigments contained within the organic film provide an active resistance against corrosion, and that the organic primer has sufficient adhesion to resist delamination from the metal surface. In terms of offering this active protection, hexavalent chromium-containing pigments carried within primer coatings have demonstrated exceptional performance in extending the longevity of corrosion protection. Chromate pigments are relatively low cost and are currently used extensively in inhibitive primers [5,6] especially due to their outstanding anti-corrosion properties for many metals and alloys [7,8]. However, the health issues associated with the use of chromates have become increasingly highlighted in recent years [9–11] although even in the late 1980s they were known to be highly toxic and carcinogenic [12]. At the same time, red lead (Pb3O4), another corrosion inhibiting species, was identified as being toxic and limitations or outright bans on the use of effective anti-corrosive pigments such as chromates and red lead continue to drive the development of less toxic alternatives [14]. Consequently, few coating types have been as heavily impacted by regulations and other non-technical agencies as inhibitive primers [13,14]. At the same time, inhibitor development is also affected by development of new resin technologies and paint formulations [14]. The composition of alkyd and red epoxy primers is discussed in detail by Hare [13], illustrating the difficulties in entering this field cold. Each primer contains over a dozen separate ingredients, from eight or 10 separate manufacturers. Hare’s [13] review provides an overview of corrosion and the principal methods of corrosion control and the design and application of primer additives. Among the many factors that makes this paper useful is that it is written from an industrial point of view and explains not only the history of additive use, and the usefulness of various chemical classes but also how changing legislation and on-going development work has impacted the field. The high toxicity and recent very strict and international legislative restrictions have led to an increasing need for the development of new pigments with reduced ecological footprints and at least similar anticorrosion properties as those shown by chromates. But due to the high number of applications and good performance of chromates, the search for new inhibitors has become a very difficult task. For instance, the more soluble chromates (e.g. zinc yellow) are currently used in alkyds, epoxy esters, and oil-modified phenolic vehicles [13]. Strontium chromate is used in epoxies for aircraft and coil coatings as well as in polyesters and urethanes [13]. At the same time, strontium chromate and basic lead silicochromate can also be used in latex systems, while basic zinc chromate is more insoluble than strontium chromate and is used in wash primers where more soluble analogues would overly decrease reaction time, film flexibility and recoat times [13]. Therefore, the performance demands for a replacement are high. All this makes the search for chromate replacement one of the drivers in changes in protective coatings, but, in the context of this chapter, is also a driver for the development of self-healing coatings. A second driver for self-healing materials and coatings development is the rising cost of maintenance for a whole range of structures. Specifically, maintenance costs of painted structures for a diverse range of applications from pipelines to infrastructure and airframes have been driving the development of structural health monitoring research [15–17] with the ultimate aim of reducing maintenance costs. Running in parallel with this research are activities on incorporating sensors into/under paint

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systems [18] as well as giving the paint systems the ability to repair, i.e. self-healing coatings [19,20]. This chapter describes the dual aim of development of multifunctional inhibitors to replace chromate as well as self-healing primer paint coatings. Ideally, self-healing implies the total recovery of the functions that the coating possessed before damage: this is total recovery of the protection against corrosion, both active and passive. Nevertheless, in practice the main characteristic of a self-healing coating is the ability to (partially) re-establish the functionality of corrosion protection of an organically coated metal. The restoration of corrosion protection can then be established by selfhealing of the polymer overlayer itself and/or the (autonomous) release of corrosion inhibitors from the organic coating resulting in reduced cathodic and/or anodic activity in the coating defect area. Therefore, the protective functions of a primer listed above, could be modified by replacing function number 4 using multifunctional inhibitors (new function 4), and the addition of another protective mechanism that is the self-healing damage repair (function 5): 4. use of multifunctional inhibitive pigments instead of ‘monofunctional’ 5. self-healing damage repair. Thus, this chapter focuses on the concept of inhibitive self-healing coatings and the selection of multifunctional inhibitive pigments to replace chromates, the rare earth metal (REM) compounds being among the most acceptable alternatives and as such is studied in more detail in this chapter. A focus is placed on the development of novel high-throughput experimental tools for identifying and evaluating eco-friendly corrosion inhibitor candidates and methodologies for their incorporation into an organic coating matrix. In addition, a particular emphasis is placed on developing primer systems for aluminium due to the high performance requirements in the aerospace industry where some of the most corrosion prone aluminium alloys are used. 9.2

Scope

For primers with self-healing and corrosion inhibitive properties, it is desirable to have one component to inhibit corrosion of the underlying metal and a second component to repair the polymer coating. These individual components can be multifunctional or multistep. For the protection of metals, many paint systems consist of an inhibitive primer and a topcoat. In case the coating is mechanically damaged or local delaminations are present, the repair function is usually achieved via release of the inhibitors from the inhibitive primer, whereby the corrosion performance of the underlying metal surface benefits from autonomous surface recovery processes by the precipitation of stable inhibitor compounds at cathodic and/or anodic sites. There are numerous challenges in the development process of novel eco-friendly inhibitors, which can be divided into six key research stages: • • • •

Development of novel eco-friendly (multifunctional) corrosion inhibitors. Evaluation of their intrinsic anticorrosion performance. Incorporation in the organic coating. Release mechanism from the organic coating and control of its kinetics.

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9.1 Inhibitor release from coating (dark grey): Creation of depletion zone (light grey) and deposition of corrosion inhibitor (central region)

• •

Investigation of the self-healing (deposition) mechanism on the corroded surface. Modification of delivery systems.

Previous studies have shown that when mechanical damage occurs, the release of inhibitor from the primer leaves an inhibitor depletion zone within the primer [21–24] as depicted in Fig. 9.1. A large dose of inhibitor is released initially but low levels of inhibitor may be released over long periods of time even up to a few years [23]. Nevertheless, the large dose of chromate is, essentially, a one-off process and not a multiple repair process. Given these factors, a new type of protecting organic coating with self-healing properties, triggered release and replenishment of the depletion zone can be proposed: a. Triggered release of inhibitor to achieve multiple dosing of the corrosion site. The ability to obtain triggered release involves the design of the inhibitor system as well as its delivery system, the latter of which may provide the trigger mechanism. In Fig. 9.2a, an example is shown of this mechanism by encapsulation of the pigments. b. Self-repair of the polymer coating, which includes repair of the defect edges to facilitate replenishment of the inhibitor at the former defect site. After the inhibitor has been released, recovery of the barrier properties of the organic coating is obtained as shown in Fig. 9.2b. c. Replenishment of the inhibitor in the depletion zone. This function will be inextricably keyed into the self-repair mechanism of the polymer coating so that the new pigment distribution will allow further release of the inhibitor if mechanical damage occurs again (Fig. 9.2c). 9.3 Development of eco-friendly inhibitors The following section describes in more detail what inhibitors are required to do, the advantages of multifunctional inhibition, and reviews the current status in the search for new eco-friendly inhibitors.

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9.2 Triggered inhibitor release and substrate protection (a); self-healing organic coating repair (b); and inhibitor replenishment of healed coating (c)

9.3.1

REM-compounds as chromate alternative

The need to replace chromate in primers on aluminium has been accepted for many years. As an example, even in the 1940s, there was an interest in organic compounds as corrosion inhibitors [25]. In 1969–70, under contract to the UK Ministry of Defence, the Paint Research Association screened some 150 compounds for their inhibitive properties. The compounds tested included known inhibitors, simple salts capable of forming insoluble aluminium salts, organic dyestuffs known to have an affinity for aluminium oxide and materials capable of complexing copper (a common alloying component in high-strength aluminium alloys enhancing corrosion [26]). This study found some 10 materials suitable for further study, but the most promising of these was incompatible with epoxy and none gave results comparable to chromate. In 1979, another 50 materials were screened, using weight loss after 3 weeks of immersion. The materials were mainly sodium salts, including nitrate, molybdate, phosphate and tungstate. Zn/Ca/Sr thioglycollate gave good results, and follow-up studies on related materials indicated that Zn thioglycollate was the most promising, but again it interfered with the drying mechanism of epoxy base resin. Thus the search

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for new inhibitors is ongoing [27–31]. Due to the existence of numerous possible inhibitor combinations, large-scale screening programmes have also emerged using high throughput methods [32–34] which will be discussed in more detail later on in this chapter. In the mid-1980s, research started on the effectiveness of rare earth metal (REM) salts as cathodic corrosion inhibitors [35]. As a result, the first papers appeared claiming that the REM ions were effective for corrosion mitigation of different metals such as aluminium alloys [35–40], mild steel [41] and zinc [42]. As an example of the interest in rare earth inhibitors in the fight against corrosion, the activity in this area can be qualitatively adjudged by just searching in an article database the keywords ‘rare earth’ and ‘corrosion’ published in this decade and this will find more than 450 articles, while in the previous decade the number is lower than 250, and in the decade of the 1980s is less than 100. So even though these numbers are only estimative and include other applications of REM as well as corrosion inhibition, they indicate that the importance of the REM is being doubled decade by decade. Thus, due to their high potential and low hazardous properties, much work has been done, especially in the last decade, in the search for new REM compounds able to provide corrosion protection equivalent to or better than chromates for a wide variety of metal substrates [35–39,43–54]. Moreover, rare earths have been modified and combined, to be used as deoxidants [55], applied in the form of conversion coatings [37,56–73], inside sol–gel coatings [74–76] or traditional organic coatings (primers) [76], in solution, and applied to many different alloys, metals and structures. For aluminium alloys, studies have been reported on AA2014 [77], AA2024 [50,51,78], AA5083 [79,80], AA6061 [81], AA7075 and AA8090 [82,83]. As inhibitors, rare earth complexes that have been tested include sulphates [81,84], nitrates [84], diphenyl phosphates [78], dibutyl phosphates [50,51] and salicylates [52,85]. Several studies have found that combinations of rare earths can produce synergistic effects. For example, Aballe et al. [80] and Davo and de Damborenea [82] found that Ce–La mixtures improved inhibitor activity. Markley and colleagues [78,86] also found that a mixture of rare earths gave results at least as good as the individual materials. Recent work has investigated whether synergies exist between the various rare earth cations using combinatorial experimental approaches [87]. Complexes such as cerium salicylate (Ce(sal)3), rare earth diphenylphosphates (RE(dpp)3) and cerium dibutylphosphate (Ce(dbp)3) have been shown to perform well in neutral aqueous solutions as well as in organic coatings [76]. Through these complexes, REM-organic compounds have shown a synergistic effect in corrosion, having both the organic ligand as an anodic inhibitor and the rare earth as the cathodic inhibitor. These studies have shown that the combination of organic– inorganic compounds can be one of the best alternatives to replace chromates due to their combined synergistic properties. Nevertheless, the mechanism of protection provided by these compounds is still under research as discussed below. 9.3.2

‘Green’ inhibitor design

For the design of a corrosion inhibitor, it is important to have an understanding of the way current systems behave. Corrosion protection, by paint systems, is currently achieved by passive (barrier coating) and active (inhibitor) protection [5,23,88,89]. Due to its inherent permeability, an intact coating allows the ingress of water into the

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paint system and dissolution of the inhibitor: at 100% RH, the water contents for coatings [44] range from 0.1–3 vol.%. If a local delamination or/and mechanical defect are also present, a release of the inhibitor into the defect area will reduce the corrosion rate. Clearly, under this scenario, the solubility and transport properties within the primer are key parameters for both inhibitor design and delivery. Over the last few decades, many studies have been performed to review, describe, elucidate and develop novel inhibitor technology [1,43–48,90–94]. It is considered that the action of inhibitors in an organic coating behaves in a similar manner to those in an aqueous environment [44]. In this context, the technical requirements of an ideal corrosion inhibitor are listed by Leidheiser [44] as follows: • • • • • •

A low solubility which exceeds the concentration required for activity. Effective in the range of pH 4–10, and preferably pH 2–12. Ability to react with the metal surface to produce a product with a lower solubility than the unreacted inhibitor. Reaction product must not reduce the adhesion of the coating. Preferably effective as both an anodic and a cathodic inhibitor. Effective against two important cathodic reactions (oxygen reduction and hydrogen production).

The mechanisms for controlling corrosion include [13]: • • • •

Resistance inhibition (decreasing the corrosion current between anode and cathode). Oxygen deprivation (depriving the cathode of fuel (e.g. barrier films) and affecting both cathodic and anodic corrosion processes). Cathodic protection (overriding all local cell corrosion activity by applying a film of an active metal, and rendering the bulk uniformly cathodic). Anodic or cathodic passivation (introducing chemicals which change the natural oxide to make it more protective and less active). In cathodic inhibition, inhibitors predominantly form insoluble precipitates or salts due to pH rise during corrosion processes [14]. Anodically active materials [92] promote the adsorption of oxygen at the surface (e.g. chromates, molybdate), or by forming insoluble complex salts with metal ions at the anodic defect sites (e.g. phosphate, borate), known as pore plugging [14].

Sinko [5] proposed that the activity of inhibitor pigments can be qualified by the inhibitor activity parameter: Ii = n

csat ccrt

[9.1]

where n is a stoichiometric factor, csat is the saturation concentration of the inhibitor and ccrt is the critical concentration of the inhibitor. As can be seen in Fig. 9.3, eventually the release of inhibitor falls below the critical concentration required to prevent inhibition and the inhibitor becomes unable to prevent further corrosion. In practice, this type of release is observed for chromate where the critical concentration is near 10 5 M [95–99]. According to Sinko [5], I should be between 1 and 100, but many inhibitors do not fall in this range. The chromate inhibitor values for SrCrO4, CaCrO4 and BaCrO4 are 5, 141 and 0.2, respectively. The rare earth inhibitor complexes are generally very insoluble and the critical inhibitor concentration generally unknown. In some cases,

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9.3 Release vs. time for an inhibitor-containing coating system – depletion of the active component (inhibitor)

this may mean that the critical inhibitor concentration may be above the solubility limit, in which case the inhibitor will never work. pH is also a variable in corrosion processes so the solubility of the compound also needs to be considered over a wide pH range. In this sense, some of the rare earth compounds that have been tested by different groups include cerium dibutyl phosphate and cerium iodate [50,100,101]. At the other end of the solubility scale are the highly soluble compounds such as the simple salts. In some respect, these materials provide different opportunities as they can be encapsulated before being combined in a paint film. In the case of the rupture of the capsule delivery system, a high initial dose will be supplied to address corrosion but no ongoing dose will be delivered. The critical concentration can be determined using standard electrochemical or exposure experiments as well as high throughput techniques. An example of the use of the high throughput approach can be seen in Fig. 9.4, where a multielectrode system (see Section 9.4 ‘High throughput screening techniques for inhibitor selection’) has been used to determine the corrosion current on a number of metal electrodes. Figure 9.4 shows the response for AA2024-T3 as well as mild steel. It can be seen that the corrosion current decreases with increasing CeCl3 concentration up to around 10 3 M which can be taken as close to the critical concentration for corrosion inhibition for AA2024-T3 and mild steel. As mentioned above, the release rate of very soluble compounds might be controlled via a delivery system. Encapsulation has been proposed as one type of delivery system [102]. This type of system tends to be one-off repair, particularly if the trigger is dissolution of the capsule due to the presence of water or fracture of the capsule due

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9.4 Corrosion current vs. CeCl3 concentration for AA2024-T3 and mild steel electrodes

to mechanical damage. It also displays problems with mixing during formulation. Other approaches for delivery systems are the use of high surface area materials, such as nanoparticles [103], layered materials, such as hydrotalcites [104–106], or mesoporous materials [107]. In this category, there are several avenues for the delivery mechanism: • •

For sparingly soluble materials, ultramilling can produce nanoparticulate powders that have a different solubility from the bulk [108]. Alternatively the inhibitor can be adsorbed onto the surface of nanoparticles or into layers in layered materials such as clays or hydrotalcites or pores of porous materials.

For materials with larger pores, it might be necessary to ‘gate’ the pores with chemicals that respond to triggers such as a change in pH. This type of approach is used with the delivery of pharmaceuticals [106,109]. Finally, changing the porosity of the carrier materials itself offers opportunities for modification of the transport properties [110]. Another key aspect of inhibitor design is the marriage of a number of different functionalities into the one complex. Some basic functionalities include: • • • • •

Cathodic inhibition Anodic inhibition Bio-inhibition Compatibility with the polymer matrix Compatibility with the healed overcoating/or the ability to form a healing overcoating

As one moves down the list incorporating additional functionalities, the challenge becomes increasingly greater and probably not linear. Therefore, for the rare earths, at least, most research has focused on a dual functionality of cathodic and anodic

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Table 9.1 Half maximal effective concentration (EC50) for SrCrO4, Ce(dbp)3 and Ce(dpp)3 (in ppm) [113]

EC50

SrCrO4

Ce(dbp)3

Ce(dpp)3

3.9

34

5.2

inhibition. Before discussing the dual functionality, it is probably worth mentioning bio-inhibition as this is often tied to the nature of the organic ligand with these rare earth complexes. There are applications where bio-inhibition is desirable such as for mould resistance in wet areas, e.g. bathrooms, or anti-biofouling such as on ships’ hulls or aircraft fuel tanks. In most instances, this will require an inhibitor that is not as environmentally friendly as desired in a chromate replacement. The rare earths themselves are relatively innocuous, so the ‘greenness’ of a rare earth inhibiting complex depends on the organic anion. An example of the ‘greenness’ of rare earth complexes, by comparison to chromate is given in Table 9.1 based on the MA-100: microtox method which is a standard aquatic toxicity test using the marine bacterium Vibrio fischeri [111]. The Ce(dbp)3 is considerably ‘greener’ since it was non-toxic whereas the Ce(dpp)3 due primarily to the phenyl groups in the organic ligand is only marginally better than SrCrO4. This is a mixed blessing since bulkier organic ligands appear to provide better protection than smaller ligands, probably due to steric effects. It follows that the relative ease of preparation of these compounds means that it is possible to modify the bioactivity through ligand adjustment. Blin et al. [112], have already demonstrated the value of this type of ligand modification by making rare earth cinnamates more soluble through substitution with hydroxyl groups to produce 4 hydroxy-rare-earth compounds. 9.4

High-throughput screening techniques for inhibitor selection

There are several experimental techniques that may be used to test the performance of a candidate corrosion inhibitor. Electrochemical testing (e.g. AC impedance and DC polarisation) are the traditional and primary research techniques, whereas weight loss determinations (immersion [45], outdoor exposure, or salt fog/cyclic chamber tests) tend to be used in commercial applications [90]. Due to the reliable and efficient performance of chromate-based inhibiting technologies, attempts to find a ‘direct replacement’ have largely failed, and therefore, present day approaches aim to develop ‘multifunctional’ corrosion inhibitors, which have the ability to passivate both anodic and cathodic reaction sites at the metal surface. Inhibitor development and evaluation approaches therefore require a large number of experiments if a systematic survey of potential cathodic, anodic and combined (multifunctional) inhibitors is to be performed [51,78]. The so-called traditional techniques mentioned above require testing times ranging from hours (in the case of polarisation studies) to weeks (for accelerated laboratory testing, mass loss, pit depth analysis, etc.) or even years (in the case of outdoor exposure testing) and produce data on only one metal type at a time. Hence, if a broad range of potential inhibitors is to be investigated, more rapid means of screening need to be found. This section therefore describes the development of rapid screening techniques for high throughput testing of inhibitors.

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Electrochemically determined corrosion rates have always been viewed as an efficient means to collect inhibition performance data (in comparison to field trials), chiefly through the estimation of a corrosion current, Icorr. With the view to creating a relatively non-destructive and more timely test method, a series of methodologies were developed in the late 1970s to enable the calculation of the corrosion current based upon linear polarisation measurements, which are recorded within approximately ±20 mV of the open-circuit potential (OCP) and therefore minimise changes to the surface arising from high overpotentials. To estimate Icorr one requires both the linear polarisation resistance, Rp and knowledge of the Tafel slopes for both anodic and cathodic processes. Mansfeld [114] developed a graphical-based method to establish Rp and Tafel slopes, and Barnartt [115,116] and LeRoy [117] utilised equi-spaced overpotentials as a basis for a simplified calculation of Tafel slopes. Bandy [118] presented a method to estimate the corrosion rate without the direct determination of Tafel constants (although they could be further deduced if desired). Errors from these approaches were up to 60%. Later, Kendig and Mansfeld [119] presented a methodology to determine corrosion rates from impedance measurements with an automatic analysis routine. Whilst these methodologies offer an ability to calculate corrosion currents, the main criticism of such methodologies is the increased uncertainty that exists for measured currents recorded close to the OCP, due to a decreased signal-to-noise ratio. This is particularly true when studying highly polarisable interfaces. In more recent times, Taylor and colleagues [32,120] have demonstrated the use of multi-electrode arrays as a means to carry out high-throughput testing on corrosion inhibitors, putting identical pairs of AA2024-T3 wires into a large number of separate reaction wells, where one electrode in each cell was polarised to 100 mV (to 425 mVSHE) with respect to a second electrode held at 525 mVSHE, corresponding to the open-circuit potential of the control [120]. The current between each pair of electrodes over time was then used as a measure of the polarisation resistance. In this manner, it was shown that large numbers of solutions could be tested simultaneously. An adaptation of the technique used by Taylor and Chambers has also been developed by our group (Muster and co-workers [34]) with several key differences: 1. 2. 3. 4.

several metals are investigated simultaneously in one solution experimental time is reduced a blank sample is used as a control for each inhibitor evaluation, and no reference electrode is used to fix electrode potentials of the metals.

Electrochemical methods are not usually carried out in the absence of a reference electrode. However, for the rapid screening of multifunctional corrosion inhibitors, the omission of the reference electrode offers the benefit of simplicity whilst retaining the key information regarding inhibition. Figure 9.5 demonstrates the current– potential relationship upon the application of a potential across two electrodes. The solid lines C1 and A1 represent the relationship of current (I) vs. potential (E) for a metal as measured by potentiodynamic polarisation experiments. C1 represents a cathodic polarisation and A1 represents an anodic polarisation. The open-circuit potential (OCP) is the unpolarised potential where the rate of cathodic surface reactions balances the rate of anodic surface reactions (i.e. where current is on average zero). By applying a 100 mV potential between two identical electrodes, one is polarised to be anodic and the other cathodic. The current flow between the electrodes

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9.5 Schematic diagram presenting the relationship between current exchange and applied potential

is determined by the stationary oxidation–reduction reactions at the anodic and cathodic electrode surfaces, respectively. The solid horizontal line connecting curves C1 and A1 represents a 100 mV applied potential and its position on the abscissa determines the magnitude of current flow. Where a cathodic corrosion inhibitor is introduced, the cathodic reactions are limited (curve C2). The dotted grey horizontal line shows that a reduction in the measured current is expected upon application of a 100 mV potential between two identical electrodes. In a similar fashion, an anodic corrosion inhibitor would decrease the measured current through a lowering of the anodic arm of the curve (not shown here). It is also noted that the position of the OCP shifts in accordance with changes in anodic or cathodic reaction rates. Measuring the current between the two identical electrodes without a reference electrode thus allows the simultaneous determination of both anodic and cathodic inhibition, which is critical for the rapid screening of multi-functional inhibitor formulations. Previous studies on corrosion inhibition have shown efficiencies to be dependent upon the concentration of inhibitor [93]. For this reason, a standard approach was devised where a series of potentiostatic scans were performed with an incremental dosing of the inhibitor. In this fashion, an indication of the ccrt, the critical inhibitor concentration, could be established. Figure 9.6 demonstrates this approach, showing the current (density) of various metals in 0.1 M NaCl solution without cerium chloride inhibitor and then at three different inhibitor concentrations (increasing from left to right). For the case shown, the current densities of the pure Mg electrode pairs were initially high and continued to increase with each step addition, indicating that cerium chloride addition is encouraging the corrosion of Mg rather than inhibiting it. On the other hand, cerium chloride acts as an inhibitor for pure zinc, AA2024-T3 and AA7075-T6, although it did not appear to have a large influence

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9.6 Current density between electrode pairs held at a potential difference of 100 mV in 0.1 M NaCl with varying additions of CeCl3. The solution was continuously purged with synthetic air and agitated using a magnetic stirrer operating at 150 rpm

on the current densities for mild steel. For the case of the two aluminium alloys, some inhibition was evident at 10 4 M, however, the data from Fig. 9.6 suggests that maximum inhibition occurs when concentrations approach 10 3 M. Zinc inhibition was also shown to occur with cerium chloride additions exceeding 10 3 M. In the inhibitor selection process, the parameter most used to classify the inhibitors is the efficiency, where the higher the efficiency, the higher the protection. Corrosion efficiencies are usually presented as a percentage according to equation 9.2: ÊI - I inhibited ˆ Eff = Á uninhibited ˜¯ x 100% I uninhibited Ë

[9.2]

where Iuninhibited and Iinhibited are the corrosion currents estimated for a metal in an electrolyte without and with an inhibitor added, respectively. Figures 9.7 and 9.8 show the ability of rapid screening inhibition efficiencies (EffME) to reflect those obtained from potentiodynamic analysis (Effpotentiodynamic) where each point represents one inhibitor. For both metals, correlations at lower ‘inhibitor’ concentrations (10 4 M) showed a significant deviation from a one-to-one relationship (dashed diagonal line). However, a good correlation in the relative performance of the inhibitors was obtained for additions of 10 2 M. It was also noted that the measured efficiencies tended to be higher than those predicted from potentiodynamic data. The differences in the efficiencies determined by the two techniques at lower inhibitor concentrations (Fig. 9.8) may be attributed to as-yet unexplored variations in the experimental design, such as exposure times to solution and inhibitor, or differences in the polarisation of the electrodes. There are also errors associated with the estimation of Tafel slopes and corrosion currents at lower concentrations, particularly where current exchanges are a mix of activation-controlled and diffusion-limited processes [121].

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9.7 Inhibition efficiencies of five inhibitor candidates at 10 2 M on AA2024-T3 (squares) and mild steel (circles) measured using the rapid screening technique plotted against the inhibition efficiencies obtained from potentiodynamic curves

In addition to electrochemical techniques that aim to determine the corrosion current, or potentiostatic current as above, high throughput studies of corrosion inhibition have been trialled for aluminium alloys with a focus on the detection of chemical changes. Chambers and Taylor [32,122] utilised two fluorimetric probes (lumogallion and morin) to complex with aluminium ions in solution. An estimate of the concentration of aluminium ions in solution was determined spectroscopically using a plate reader. In this way, 14 inhibitors were investigated in terms of their synergistic and antagonistic behaviour upon mixing. With a similar motivation, the same authors also presented a novel 96-electrode determination of surface copper enrichment on aluminium alloy wires after exposure to various inhibitor solutions using a micromultielectrode analyser [123]. This methodology enabled 11 inorganic inhibitors and 44 binary combinations at seven different concentrations to be determined. Combinations of sodium molybdite paired with cerium, yttrium or lanthanum chloride were found to rival the performance of sodium chromate in minimising the presence of surface copper. Another novel method for accelerated corrosion testing developed by our group has been reported by White et al. [124], which enables high throughput corrosion experiments based on a microfluidics design. The experimental rig is depicted in Fig. 9.9. Low profile channels (200 μm) in height were cast into PDMS (polydimethyl siloxane). The negative pattern of the channel design was created using a UV light curable flexoplate sheet. The channels were then cast onto the patterned flexoplate positive master. Finally, inlet and outlet holes were cut through the PDMS. The

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9.8 Inhibition efficiencies of five inhibitor candidates at 10 4 M on AA2024-T3 (filled squares) and mild steel (red circles) measured using the rapid screening technique plotted against the efficiencies obtained from potentiodynamic curves

9.9 High throughput corrosion experiment based on Microfluidics

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solution was pumped from the bottom of the channels to the top of the channels after it was clamped to the substrate that was to be tested. There are many different types of experiments that can be performed in this set-up including: •

•

•

•

Immersion testing: In this type of experiment, NaCl solution with inhibitor is pumped through the channels for a period (typically 24 h). The solution and the corrosion product left on the surface are then analysed and the degree of corrosion is determined. This will be discussed in more detail below. Filiform corrosion: With rolled surfaces without any further treatment, filiform corrosion occurs during a typical 24 h experiment. When an inhibitor is present, filiform may or may not occur within this period. This is a very rapid technique to determine the resistance to filiform corrosion. Crevice corrosion: the PDMS channel mask is permeable to O2, however if that supply is cut off, by metallising the surface or enclosing the experiment in an inert atmosphere, then the only source of O2 is through the circulating solution. This can be depleted of O2 to simulate a crevice. Leaching from primers: The multichannel array is not limited to corrosion experiments and can be used to examine the release of inhibitor from a primer under a range of conditions e.g. pH.

The corrosion experiment, however, is the focus of the rest of this section which aims to describe the experiment itself as well as demonstrate the degree of acceleration over standard immersion experiments. In a typical test with this set-up, the PDMS mask with the channels is clamped over the test panel and the test solution is pumped across the surface. The test solution can be pumped either as a once through experiment or backwards and forwards. After the experiment, the solution is analysed using Inductively coupled plasma atomic emission spectroscopy (ICPAES). A second PDMS master with slightly wider and longer dimensions is then placed over the original channels and a wash solution containing CrO9/H3PO4 is then pumped through the channels to dissolve the remaining corrosion product. This particular experiment is designed to be an accelerated version of a standard immersion/weight loss test. This has been validated by White et al. [124] who have compared the multichannel approach with 10 channels to a standard immersion for a number of well-researched inhibitors including chromate, cerium nitrate, cerium iodate, and cerium salicylate as listed in Table 9.2. Table 9.2

Inhibitors, concentrations and channels

Channel number A B C D E F G H J K

Inhibitor

Concentration (M)

K2Cr2O7 " Ce(NO3)3 " " Ce(IO3)3 " Ce(salicylate)3 " "

10 4 10 5 10 3 10 4 10 5 10 4 10 5 10 3 10 4 10 5

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9.10 Ten-channel experiment of 10 different inhibitor solutions after 6 days of continuous flow at 0.05 ml/min, comparing the visual average ranking (1 = best inhibition) of 10 inhibitors with the ranking assessed by Al ICP analytical ranking

Clearly, there are several ways of ranking the outcome of an experiment like this including the method described above of using ICP-AES analysis. However, visual observation is the quickest and for the purposes of screening the most convenient. Thus, to test the validity of visual ranking, the corroded and stripped metal specimens were assessed by 10 people with some degree of corrosion knowledge and given a ranking from 1 (best) to 10 (worst). The average and spread of the rankings were compared with the dissolved Al ICP-AES results which were also ranked from 1 to 10 (Fig. 9.10). The visual correlation agreed well with the ICP ranking with only two inhibitors falling greater than one ranking outside of each method. In screening inhibitors, these results indicate that visual assessment can be used for a first pass evaluation of the performance of the inhibitor. However, from a scientific viewpoint, it is also desirable to demonstrate that the corrosion processes/mechanisms are occurring as in standard testing but at an accelerated rate. This latter objective is more of a challenge. White et al. [124] have established that the degree of acceleration is approximately one order of magnitude faster than the standard immersion test. The corrosion rate as a function of the flow rate through the channels is a good place to obtain an understanding of the nature of the corrosion processes in the channel. Figure 9.11 shows the amount of dissolved Al determined using ICP-AES plotted for a 24 h experiment. It can be observed that over 24 h, there is considerable variation in the dissolved Al with flow rate. A flow rate between 0.25 and 1 ml/min gave a maximum degree of corrosion of both Al and Cu (Cu results not shown here); 1000 μg of Al was removed in 24 h at a flow rate of 0.25 ml/min corresponding to an average depth of removal of 1.6 μm for the channel. At a flow rate of 2 ml/min, however, corrosion was significantly reduced since the flow rate was demonstrated to be too high to allow for stable pitting to occur. These results provide some insight into the corrosion process in the channels and its

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9.11 Al corrosion vs. flow rate at a fixed volume (20 ml) for a fixed time (24 h)

relationship to the immersion experiment. Part of the corrosion process involves the formation of caps of corrosion products over actively corroding sites which create an acidic environment that protects a growing pit. At high flow rate this product is immediately removed downstream hindering the formation of the acidic pit solution thereby inhibiting pit growth. In earlier work, Ilevbare et al. [125] also noted that mass transport control may have an important role to play in S-phase dissolution as well as pH. At more moderate flow rates, this cap can form providing an environment for the formation of pits. The acceleration is probably provided by the increased oxygen supply due to flow of fresh solution over the surface versus static immersion conditions. This is supported by the data in Fig. 9.12 where the pit size distributions are plotted for immersion, 0.05 and 0.25 ml/min. It can be seen that the slower flow rate has a pit size distribution similar to the immersion pit size distribution, whereas the pit mouth areas for the faster flow rate are much larger. While the focus of the foregoing discussion has been on comparison with standard types of immersion experiments, this type of experiment is well suited to a combinatorial approach for inhibitor development. This can be achieved by mixing inhibitors from different reservoirs before pumping through the channels. This approach offers the opportunity to explore a large variable space in one experiment and is the subject of further research. Finally, it is important to understand the limitations of each testing method [90], and that no one method will provide all of the necessary information. Thus, it is important to confirm rapid screening techniques with more traditional and industryaccepted methodologies such as weight loss tests or salt fog spray tests. For example, zinc phosphate has been shown to perform poorly in accelerated tests but generally well in exposure testing [47]. 9.5

Inhibitor incorporation into organic coatings

Once the inhibitors have been selected, the next step is their incorporation into organic coatings. This new step again presents some difficulties and a lot of extra

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9.12 Comparison of the pit mouth area between the static immersion experiment (blue) and the continuous flow channel at 0.05 ml/min (orange) and 0.25 ml/min (yellow) of the results in Fig. 9.11

research to develop the optimal process of incorporation of the inhibitors to avoid undesirable reactions between the polymer matrix and the inhibitor. As to the incorporation of the pigments into polymer matrices, a lot of effort is being expended and some of the results and advances will be presented here. Organic coatings have a long history of providing a barrier for oxygen and moisture transport to metallic surfaces and thus preventing corrosion. However, all organic coatings are permeable and, with time, conditions become more conducive to mass transport and, in case some local delaminations or contaminants coming from the substrate or paint application are present, corrosion processes can be initiated due to water reaching the interface. In the past, chromate species have been the highest performance corrosion inhibitors available and have been used to improve the corrosion protection properties of many organic coatings. The main factors required for successful inhibition of corrosion are: 1. The inhibitor needs to have limited solubility. When incorporated into the coating matrix, the inhibitor needs to be sufficiently soluble for transport within the coating to the metal–coating interface where corrosion occurs, however, if the inhibitor is too soluble then osmotic blistering occurs or the inhibitor could be washed out of the coating system. 2. The inhibitor system must be stable during coating formulation, application and cure of the polymeric components of the coating. 3. The inhibitor must have high corrosion inhibition efficiency. There have been several strategies pursued to create organic coating systems containing chromate-free corrosion inhibitors; these include simple addition of corrosion inhibitors as pigments, microencapsulation and release from molecularly engineered

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layered materials. The inhibitors have been added to a wide range of conventional coating systems including epoxy-amine and epoxy-polyamide polymers [104, 126–131], alkyd resins [132], poly(vinyl butyrate) [105,133] and poly(phenylene sulphide) [128]. The use of corrosion inhibitors as pigments, or multifunctional filler particles, is one of the most widely considered and patented approaches. Schuman et al. [126] investigated the use of cerium complexes as corrosion inhibitors in epoxy coatings for aluminium aerospace alloys. Their research showed that cerium reacted rapidly with amine and polyamide components of the epoxy systems causing a change in colour of the system and the evolution of heat. In addition, there was a reduced storage life and perhaps altered reaction kinetics. The cerium–resin interactions did not affect the corrosion inhibition, with the cerium containing primers successfully inhibiting corrosion in accelerated testing using neutral and SO2 salt spray and filiform corrosion test methods. Other recent work has shown that cerium(III) and mischmetal diphenyl phosphate additions to epoxy primers can provide synergistic corrosion inhibition and inhibit corrosive processes [111]. Other forms of coating systems which contain corrosion inhibitors include in-situ phosphatising coatings such as those described by Neuder et al. [126] where functionalised aryl-phosphonic and arylphosphoric acids were added to an epoxy primer. The inhibitors were shown to deposit uniform metal-phosphate products on the substrate surface and provide a similar level of performance to chromate-inhibited control samples when tested over 3000 h in a neutral salt spray test. The addition of water soluble corrosion inhibitors to coating systems can be problematic due to the solubilisation of the inhibitors causing osmotic blistering. One approach is to coat inhibitors with a plasma-polymer surface to reduce the release rate of the encapsulated inhibitor. Yang and van Ooij [131] coated triazole particles with an inner layer of (hydrophobic) perfluorohexane plasma polymer and then an outer layer of pyrrole plasma polymer. The resultant particles were then dispersed into a water-based epoxy and it was shown that the release rate of the triazole had been significantly reduced and blistering was not observed. Another interesting use of inorganic pigments was the addition of hydraulic calcium aluminate fillers to poly(phenylene sulphide) coatings where they assisted polymeric coating self-repair in harsh (200°C) brine environments. When cracks were created in the polymer, a decalcification–hydration reaction of the fillers led to the dissolution of calcium bicarbonate and the formation of block-like boehmite crystals in the damage sites of the coating. The boehmite crystals grew to approximately 4 μm in size sealing the cracks and restoring the function of the coatings as a corrosion prevention barrier [134]. Whilst the majority of inhibiting pigments are simply added or milled into the organic base resin, there is a significant body of work that has investigated the release of inhibiting pigments from ion-exchange structures. The use of natural and synthetic ion-exchange structures such as hydrotalcites and clay minerals has been investigated widely and shown to be a promising approach to deliver both anions and cations to suppress corrosion. The layered structures act as reservoirs for the corrosion inhibitors and allow the release of the inhibitors when triggered by moisture and through pH changes caused by corrosion reactions. The ion-exchange structure may support anionic or cationic exchange. Williams and McMurray [133] showed that release of organic anions from hydrotalcite pigments in poly(vinyl butyrate) coatings could inhibit filiform corrosion with

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corrosion inhibitive efficiency rankings decreasing in the order benzotriazolate > oxalate >> ethyl xanthate. Buchheit and coworkers have extensively studied the use of Al–Zn–vanadate containing hydrotalcites as both pigments [104] and components for conversion coatings [136–138]. When incorporated into epoxy resins, the vanadates acted as an anodic inhibitor and the Zn2+ as a cathodic inhibitor. They also showed that it was possible to use the pigments as a sensor of the remaining service life by monitoring the characteristic changes in the X-ray diffraction patterns of the pigments which provided information on the relative amounts of chloride and vanadate complexed to the hydrotalcite. This technique would be able to determine when the pigment had released all of the inhibitor species and therefore provide a means for determining when the coating system needed replacement. Where cations such as lanthanides, zinc and calcium are able to be released, they have demonstrated an ability to inhibit corrosion. Böhm et al. [139] examined the release of Ce3+ and Ca2+ from bentonite pigments; they showed that the calciumexchanged pigment matched the anti-delamination performance of a strontium chromate containing control, and the cerium containing pigment provided superior anti-delamination performance compared to the chromate containing system. In summary, a number of successful methods have been used to incorporate non-chromated inhibitors into organic coating systems. A wide range of inhibition systems have potential as chromate replacements with low solubilities, which do not cause osmotic blistering but with sufficient transport to allow movement to corrosion sites. No long-term performance testing has been reported for these approaches, and further research has still to be done to find the proper system(s) to replace chromates, hence the viability of these newer systems over the service life of a structure (2–25 years) is largely unknown 9.6

Release kinetics

Apart from the incorporation of the inhibitors to the organic matrix, another important factor that has to be controlled is the release kinetics of the incorporated pigment and its optimisation. The release of dispersed inhibitors in an organic coating is a three-stage process. Firstly, there needs to be the ingress of a triggering substance, usually water, through the coating to the dispersed inhibitor particles. Typically, this will occur because of a breach or other kind of defect in the overlying protective top coat or through general permeation of a primer. Secondly, there needs to be a reaction, usually dissolution, on the surface of the inhibitor particle which releases the active inhibitor agent; and finally, this inhibiting agent must then move to the site where it can perform its protective action. Straightforward inhibitor particle dissolution is the most common release mechanism, and may of course be influenced by water, pH, presence of other ionic species, or further speciation reactions. It is however a rather crude means of providing control over dosage rates, since it is also dependent on temperature and common ion effects [4]. Other possible release mechanisms include hydrolysis reactions and ion exchange [4]. Release by hydrolysis or acid/base reactions offers a degree of response to the corrosive environment, although generally to corrosive factors rather than initiators of corrosion such as chloride and sulphate [4]. Ion exchange would seem to be, in theory at least, the most promising process as it could limit the amount of electrolyte present and prevent any corrosion reaction taking place [4]. Nevertheless, this mechanism has an important drawback because the surface area of active

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pigments is too low to act as a major release mechanism, with the possible exception of red lead (Pb3O4) [4]. It should also be applicable to use low levels of otherwise highly-soluble inhibitors [4], although the only material in commercial use to date is calcium-exchanged silica, even though the action of this material is not only due to the presence of the calcium, but is also affected by the support itself [4]. It is tempting to assume that the saturated solubility of the inhibitor in bulk solution should be a good indicator of its release kinetics. However, although this solubility is a factor, it is not always a good predictor of the overall release kinetics as measured by leaching studies [140,141]. Indeed, Nazarov et al. [141] have shown that for coatings nominally identical, except for the inhibitor used, the trends in inhibitor leaching rate are not necessarily well correlated to the bulk solution solubility. This is attributed to possible physical and physico-chemical interactions between the inhibitor and various functional groups in the polymer matrix and the other fillers or pigments present in the coating. Moreover, the chemical environment in the pore space of a heterogeneous medium could conceivably be quite different from that of a bulk solution. Such interactions can occur during each of the stages mentioned above – ingress, dissolution and subsequent transport of the inhibitor species. Other factors which need to be considered are any mechanical and structural changes in the coating that occur during leaching or over longer periods of time. Such changes will influence the release kinetics by modifying the available molecular transport pathways. These effects can differ appreciably between different coatings [142,143]. The overall picture that emerges is that often each inhibitor–polymer combination may be quite unique. The time profiles of the release are also of practical interest as it is desirable to have a controlled release – ideally a high initial release dose to provide protection, followed by a tapering-off to avoid unnecessary wastage of inhibitor. Some studies show time profiles for leaching of chromate pigments and a variety of chromate replacement inhibitors [141,144]. It was found that chromate and vanadate exhibited a logarithmic, or even slower than logarithmic, growth in the cumulative amount of inhibitor leached out with time. This corresponds to a kind of self-quenching behaviour, with between 24% and 45% of the total amount released over 15 days being released in the first hour after water is introduced. This kind of behaviour cannot be explained by a simple diffusion model, which would predict an inhibitor release with an approximate t1/2 behaviour with time. Other studies have also observed non-t1/2 profiles [22,23]. On the other hand, Nazarov [141] studied other inhibitors, phosphate and tungstate, which showed approximate t1/2 and t behaviours, respectively. These responses can be explained on the basis of standard diffusion or dissolution being the rate determining steps. Putting aside the absolute amount of inhibitor released, these different time profiles indicate that there are different mechanisms, or at least different balances between mechanisms, at work in the various cases, and as a consequence, every pigment–organic matrix system has to be studied as a specific case. 9.7 9.7.1

Protection mechanisms REM deposition mechanisms

In previous sections, rare earth metal (REM)-based inhibitors have been acknowledged as good candidates to substitute the chromate-based inhibitors. Nevertheless, despite the numerous studies performed to understand the mechanisms of protection they offer [39,58–73,145,146], the exact mechanism of deposition is still under debate

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and moreover, seems to depend on the type of substrate that it is protecting. In this section, the supposed mechanisms by which REMs protect the substrate are analysed. Despite the exact mechanism not being clear, the process by which the REM, and Ce cations in particular, protect the substrate seems to be generally accepted and common for all of the studied metals. Thus, it is believed that the main explanation for the Ce(III) aqua cation to protect the metallic substrates against corrosion, is based on the creation of an oxy-hydroxide film on the cathodic sites of the metal. Moreover, the influence of the Ce3+ to reduce the corrosion process relies on the reduction of the cathodic currents through the barrier protection offered by the cerium hydroxide or even by the process of oxidation from Ce(III) to Ce(IV). In the corrosion process of metals, the cathodic reaction mechanism is dependent on the pH of the media. In neutral and basic media [40], the reduction reactions are: O2 + 2H2O + 4e ȥ 4OH

[9.3]

O2 + 2H2O + 2e ȥ H2O2 + 2OH

[9.4]

2H2O + 2e ȥ 2OH + H2

[9.5]

While, in acid media the main reactions are: O2 + 4H+ + 4e ȥ 2H2O

[9.6]

2H+ + 2e ȥ H2

[9.7]

where reactions 9.3, 9.4 and 9.6 occur in the presence of oxygen and reactions 9.5 and 9.7 in non-aerated solutions. For the case where localised corrosion mechanisms are limited (i.e. pitting or crevice corrosion), the main cathodic reactions are those given by reactions 9.3 and 9.4, the latter being a secondary and slower reaction. At the anodic sites, the dissolution of the metal is described by the simplified reaction in equation 9.8: M ȥ Mn+ + ne

[9.8]

In the absence of an inhibitor, the combined reactions would give rise to the formation of the dissolved metal cations whose stability depends on the pH of the media as well as the presence of aggressive ions such as Cl . Nevertheless, when cerium or other REM is present in solution, the reactions become more complex. On the metal surface, the reaction given by equation 9.9 will take place, leading to the deposition of cerium(III) hydroxide at the cathodic sites due to the local pH increase produced by reactions 9.3 and 9.4 [39,40,147]: Ce3+ + 3OH ȥ Ce(OH)3 Ȧ

[9.9]

At the same time, in solution, oxidation can occur by equation 9.10 giving rise to the formation of soluble Ce(IV)-hydroxy compounds [40,148] via: Ce3+ + 2H2O ȥ Ce(OH)22+ + 2H+ + e

[9.10]

This oxidation reaction of Ce(III) to Ce(IV) can take place if the pH is above 8.7 [40], a situation that is created near the surface of the metal due to reactions 9.3 (four-electron pathway) and 9.4 (two-electron pathway), that produce, respectively, an increase in the local pH to values of 10.8 and 10.5 [40].

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According to the Pourbaix diagram [147,148], there is a large influence of pH on cerium deposition. In aerated solutions, if the pH is higher than 8, the Ce(III) becomes oxidised to Ce(IV) at or near the metal–solution interface, and the resulting hydrated Ce(IV) oxide films are more effective barriers to ongoing cathodic activity [39,149]. Moreover, this oxidation rate process can be increased if hydrogen peroxide is present in the solution which will reduce the pH at which precipitation of cerium compounds occurs [147,150]. In the case of cerium-based conversion coating baths, added hydrogen peroxide has been proven to reduce the time of cerium oxide formation to minutes instead of days that would be necessary in normal conditions [71]. On the other hand, hydrogen peroxide can be formed in solutions where the pH is between 7 and 10 and potentials between –0.6 to –0.7 V vs. SCE [149] as shown in equation 9.4. Thus, the oxidation reaction of Ce(III) to Ce(IV) can occur via two different paths [40]. The first path would be following the four-electron cathodic reaction (equation 9.3) which will lead to: 4Ce3+ + O2 + 4OH + 2H2O ȥ 4Ce(OH)22+

[9.11]

This first oxidation reaction will take place in regions just above the surface of the metal because one oxidant is produced at the surface and has to diffuse into the solution and the other oxidant is in the solution and has to diffuse to the surface. If the two-electron reaction path takes place or hydrogen peroxide is added to the solution, then the preferential reaction will be: 2 Ce3+ + 2OH + H2O2 ȥ 2Ce(OH)22+

[9.12]

The products coming from this second oxidation reaction will be generally deposited on top of the surface because the two oxidants are produced just on the surface of the metal promoting a better and more efficient deposition of the cerium on top of the metal surface to be protected. The process initiated by equations 9.9 and 9.12, and to a minor extent reaction 9.11, will continue by reactions 9.13 to 9.15, leading to a cerium oxide/hydroxide layer formation [151,152]: Ce(OH)3 ȥ CeO2 Ȧ + H2O + H+ + e−

[9.13]

Ce(OH)22+ + 2OH ȥ CeO2 Ȧ + 2H2O

[9.14]

Ce(OH)22+ + OH + e ȥ Ce(OH)3 Ȧ

[9.15]

As a result of reactions 9.9 to 9.15, a layer of cerium oxy-hydroxide with extremely low solubility is formed on top of the surface starting at the cathodic areas and extending all over the surface by a nucleation and grain growth mechanism [152]. The general equation of the process with hydrogen peroxide will then be written as [153]: 2Ce3+ + H2O2 + 6OH− ȥ Ce(OH)4(CeO2·2H2O) Ȧ

[9.16]

Or more likely as [59]: Ce3+ + 3/2 H2O2 + 3OH ȥ Ce(O2)(OH)2 Ȧ + 2H2O

[9.17]

The peroxo complex in equation 9.17 has been identified in cerium conversion coatings and should be formed in aqueous solutions where reaction 9.4 is taking place. Moreover, as demonstrated by Lau et al. [154], a mixture of Ce2O, Ce(OH)4 and CeO2(OH)2 may be observed on the surface of cerium-based conversion coatings.

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According to Davenport et al. [39], when a piece of aluminium is immersed in a bath containing Ce3+ ions at a neutral pH, the cerium will deposit as Ce3+ with a long exposure time (5 days), probably following equation 9.9. Nevertheless, when time of exposure is increased to 7 days, the deposit having Ce3+ was oxidised to Ce4+, probably indicating that reaction 9.13 took place indicating that reactions 9.14 and 9.15 are hindered due to the effect of the cerium hydroxide layer on the cathodic reduction reactions. With respect to the influence of the substrate in the protection process, the early work clearly indicated that the rare earths protected the surface through cathodic inhibition as seen through the reduction in the current density on the cathodic arm of polarisation curves as well as in electron microscopy [35–38,153]. Solid solution alloying additions have a strong influence on the deposition kinetics on the matrix alloy as demonstrated by Hughes et al. [71] for a range of alloys. The Ce3+/Ce4+ ratio may also change with alloying additions such as Cu since Davenport et al. [39] observed higher amounts of Ce4+ when the substrate was copper instead of aluminium [39]. Thus, the influence of this element (and other alloying additions) needs to be considered when working with aluminium alloys with high copper content such as AA2024 or AA7075. Moreover, in their work with pure zinc and iron Böhm et al. [149] proposed that, to allow the deposition of cerium oxide on top of the substrate, the existence of an initial hydroxide layer of the metal itself is necessary, meaning that, in aggressive media where the initial hydroxide is rapidly removed will result in a lower deposition of the cerium compounds. Thus, when cerium complexes are introduced into organic layers and leached out from the coating, it is expected that if the cerium cation goes into solution then it will act in the same way as explained above for the immersion studies. So the protection layer will be formed by deposition of cerium oxide-hydroxides onto the cathodic areas such as cathodic intermetallic particles, and those anodic intermetallic particles that have been converted to cathodes through dealloying such as the S-phase (Al2CuMg) in AA2024. 9.7.2

REM-organic compound deposition

To this point, the discussion of the mechanism of deposition and inhibition has focused on the rare earth cation in the presence of OH , peroxo derivatives and an implied Cl concentration. One of the fundamental tenets of building a multifunctional inhibitor is designing a compound (rare earth in this case with organic anion) that can deliver its own anion to the corrosion site. Once at the corrosion site, pH excursions related to local electrochemical reactions result in the breakup and reaction of the separate components of the multifunctional inhibitor with the surface. In this sense, the mechanism of inhibiting organics in near-neutral solutions has been studied by ellipsometry, UV/vis, IR XPS, and Auger spectroscopy and electrochemical techniques [91]. Romagnoli and Vetere [47] go into some detail about the various testing carried out on zinc phosphate and how the type of test, the type of resin (e.g. water-based epoxies) and fillers (e.g. magnesium silicate or titanium oxide) used and the type of result (e.g. lack of corrosion, but significant chalking) can affect the test. Particle diameter, PVC, additive and binder all affect the outcome and are often not specified in the literature [47]. The pigment activity of an anticorrosive paint only becomes important once the barrier effect of the film has broken down [47].

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The utilisation of organic anions as components of a multifunctional inhibitor complex has largely been successful because of their ability to form barrier layers on the metal surface. For example, carboxylic acids have long been known to impart some degree of protection to metal surfaces [156]. Deposited films of polypyrrole [157] and organosilane [158] films on aluminium alloys have improved the resistance to corrosion, while adsorbed surfactants have also been successful by impeding the diffusion of corrosive agents into the matrix [159]. Lamaka et al. [160] reported that salicylaldoxime, 8-hydroxyquinoline and quinaldic acid formed thin organic layers of insoluble complexes on an AA2024 surface, which suppressed the dissolution of Mg, Al and Cu from intermetallic particles. Zheludkevich et al. [161] treated the surface of AA2024 with derivatives of triazole and thiazole and the resulting organic layers on the surface decreased the rate of both the anodic and cathodic processes. 9.7.3

Surface analysis techniques

There have been several studies that have attempted to identify the species that have reacted on the surface, and finding the techniques that have the sensitivity to detect species is a major challenge. This is particularly so for the anionic component since many of these species passivate the surface by forming a protective layer while the detection of a high atomic number species like the rare earths presents less of a problem since they have a high X-ray cross-section for nuclear and X-ray techniques. Some techniques that are useful for studying this complex problem are listed in Table 9.3. The table lists a number of techniques that cover elemental analysis to vibrational spectroscopy. It also gives an appraisal of the likelihood of detection of the different inhibitor components themselves as well as an assessment of studying the interaction of the incorporated inhibitor with the paint system – in this case an epoxy system. Finally, it should be pointed out that techniques need to be developed to demonstrate self-healing and many of the current techniques are not suitable. This should only be taken as a guide since many anionic inhibitors may produce a characteristic signal that can be gainfully used to study their interaction with the surface and accelerators and provide opportunities not available with laboratory-based equipment. Some of the challenges in design and characterisation can be demonstrated by looking at a case study: the interaction of Ce-dibutyl phosphate with AA2024-T3 [152,153]. The reason that the characterisation is important is that knowledge of where the two components of the multifunctional inhibitor are interacting with the surface feeds back into the design of the inhibitor. Ideally, it would be desirable to know where all components of the inhibitor are interacting with the surface. This dictates an approach with high spatial resolution as well as sensitivity. Thus, vibrational techniques such as Raman and FTIR are chosen for their sensitivity to molecular bonds, which is good for the organic component but poor for the inorganic components. So in Fig. 9.13, Raman vibrations for functionality from the anionic components are mapped and compared to an optical image of a feature on the surface. There is clear correspondence of the corrosion feature with intensity from bands at 459, 843 and 1073 cm 1. The phonon band from CeO2 (459 cm 1) suggests the presence of a cathodic particle at the centre of the feature. The bands at 843 and 1073 cm 1 represent a C–C skeletal and C=O symmetric stretch, respectively, suggesting that the dibutyl phosphate anion surrounds the cathodic site. These features

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Table 9.3 Techniques for detecting the reacted species on surfaces and their sensitivity Techniques Cation

Anion

Epoxypolyamide

RBS

Good sensitivity and spatial resolution Good surface sensitivity, but poor laterally Good sensitivity and moderate spatial resolution Poor

Poor sensitivity

Poor sensitivity

EDS moderate sensitivity and good spatial resolution Good sensitivity and spatial resolution

Poor sensitivity

XPS

TOF SIMS

Vibrational

SEM/EDS

PIXE

Positron Poor techniques

Incorporated inhibitor

Self-healing

Good sensitivity and spatial resolution Composition Composition and chemical and state chemical state Moderate Composition sensitivity, and moderate chemistry spatial resolution

Good sensitivity and spatial resolution Good on interactions with surfaces May be poor depending on the nature of inhibition

Good sensitivity and Raman has good spatial resolution Poor sensitivity

Mixed sensitivity and good spatial resolution for FTIR and Raman Mixed sensitivity and resolution

Moderate sensitivity good spatial resolution

Poor sensitivity

Mixed sensitivity and resolution

Poor

Possibilities with encapsulated system

Positron techniques voids – no spatial resolution

Raman FTIR moderate sensitivity good spatial resolution EDS moderate sensitivity and good spatial resolution Excellent spatial resolution and sensitivity for cationic components Poor

Poor sensitivity

Moderate sensitivity, moderate spatial resolution Good sensitivity and spatial resolution

RBS, Rutherford backscattering spectroscopy; XPS, X-ray photoelectron spectroscopy; TOF-SIMS, time of flight secondary ion mass spectrometry; Vibrational spectroscopies include Infrared techniques such as FTIR and Raman spectroscopy; SEM/EDS, scanning electron microscopy with energy dispersive X-ray emission spectroscopy; PIXE, particle-induced X-ray emission spectroscopy.

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9.13 Raman map of AA2024-T3 after 24 h exposure to 200 ppm Ce(dbp)3 in 0.1 M NaCl. Top left: 1073 cm 1 = C=O; Top right: optical micrograph; Bottom left: 459 cm 1 = CeO2 phonon peak; Bottom right: 843 cm 1 = C–C (reproduced with permission)

indicate that the anionic and cationic components of the inhibitor phase are playing their separate roles. Unfortunately, the CeO2 phonon mode only exists for ceria and so cannot be used for other rare earths. So other mapping techniques that rely on X-ray emission are preferred more generally for the rare earths. A second approach is to manufacture model compounds of intermetallic (IM) particles that appear in the alloy of interest and study the reactions of the inhibitors with the modelled IM compounds. Hinton et al. [162] and Scholes et al. [163] have studied the interaction of Ce(dbp)3 with IM compounds that are in AA2024-T3. Figure 9.14 shows XPS survey spectra of pure Al, AA2024-T3 and a number of IM compounds after exposure to salt solution containing 200 ppm Ce(dbp)3. There is little obvious difference between the survey scan on pure Al and AA2024-T3. The spectrum for the S-phase (Al2CuMg) shows some Cu and probably higher Mg. The reason for such a small difference

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9.14 XPS spectra of the surface of (a) AA2024-T3, (b) Al (99.999%), (c) Al2CuMg, (d) Al7Cu2Fe and (e) Al3Fe after 24 h immersion in 0.1 M NaCl solution containing 200 ppm Ce(dbp)3

between the S-phase and the Al alloys is that there is often surface enrichment of alloying components such as Mg and Al during a range of preparation conditions [164–168]. The S-phase is also likely to be a reaction site where the anionic inhibitor will react, and the photoionisation cross-section for P is low leading to low sensitivity [169]. On the other hand, Ce is evident in the survey spectra of Al7Cu2Fe and Al3Fe. The detailed spectra of Ce for these materials indicated that Ce3+ is present on Al7Cu2Fe but Ce4+ is present on Al3Fe. The different oxidation states for the cathodic IM particles suggest differences in the details of the reaction with the surface of the IM particles. XPS only gives information on the cationic Ce species. To obtain information on the anion species, as seen above, vibrational techniques can be used but for this particular experiment they did not provide much information on the reaction with the surface. This may in part be due to the thicker oxides that were on the surface that may have obscured the anionic inhibitor phase. In this instance, Rutherford backscattering spectroscopy (RBS) was used. RBS is a nuclear technique which relies on measuring the energy of either ź-particles or protons backscattered from nuclei in the sample. This higher atomic number species will generally have greater sensitivity.

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9.15 XPS results for FeAl3 and Cu2FeAl7 reproduced IM

Thus, in Fig. 9.15, the Ce peak for the Cu2FeAl7 is very intense, but the lighter elements show only small steps or peaks and represent a distinct inhibitor layer on the intermetallic phases that are rich in cerium but, as the XPS showed, not necessarily CeO2. P was detected in Cu2FeAl7 but at very low levels and could only be extracted using fitting techniques. FeAl3, for example, had no significant signal for Ce, but more P was detected. These results indicate that the cathodic and anionic components of the inhibitor react in different ways on different IM particles. Therefore, the deposition reactions of the inhibitor with the alloy surface must be couched in the context of the specific reactions with individual components on the surface. In conclusion, the deposition process of the cerium ions on the surface can be summarised as first, deposition of Ce(OH)3 and later, a reaction with oxidants such as the cathodically generated H2O2 producing hydrated CeO2, with a middle step of Ce(OH)22+ production. The protection offered by the REM oxide films (mainly Ce(OH)3 and CeO2) is based on the electrical insulator capability of those oxides, which will allow the reduction of the cathodic reactions, together with their barrier properties as an oxide layer. On the other hand, the protection offered by the anionic organic component is mainly through formation of a layer on top of anodic sites increasing the barrier properties of the previously exposed surface. Acknowledgements The authors would like to acknowledge P.A. White, D. Lau, T.G. Harvey, P.A. Corrigan and S.G. Hardin for their very valuable contributions to this work. The authors also gratefully acknowledge the financial support from the Innovative Research Programme in The Netherlands (IOP-Innovatiegerichte Onderzoeksprogramma’s) for this research under project number IOPSHM0633. References 1. A. D. Mercer, Spec. Publ. R. Soc. Chem., 71 (Chem. Inhib. Corros. Control), (1990), 46–56.

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122. B. D. Chambers and S. R. Taylor, Corros. Sci., 49 (2007), 1597. 123. B. D. Chambers and S. R. Taylor, Corrosion, 63(3) (2007), 268–276. 124. P. A. White, A. E. Hughes, S. A. Furman, N. Sherman, D. Lau, T. H. Muster, P. A. Corrigan, M. A. Glenn, T. G. Harvey, S. G. Hardin, J. Mardel, S. J. Garcia-Espallargas, C. Kwakernaak and J. M. C. Mol, Submitted 2008. 125. G. O. Ilevbare, O. Schneider, R. G. Kelly and J. R. Scully, J. Electrochem. Soc., 151 (2004), B453. 126. H. Neuder, C. Sizemore, M. Kolody, R. Chiang and C. T. Lin, Prog. Org. Coat., 47 (2003), 225. 127. V. Sauvant-Moynot, S. Gonzalez and J. Kittel, Prog. Org. Coat., 63 (2008), 307. 128. T. Sugama and K. Gawlik, Mater. Lett., 57 (2003), 4282. 129. N. Voevodin, D. Buhrmaster, V. Balbyshev, A. Khramov, J. Johnson and R. Mantz, Mater. Perf., 45 (2006), 48. 130. T. S. Schuman, A. Stoffer J.O., in International Waterborne, High Solids, and Powder Coatings Symposium, New Orleans, LA, USA, 2002. 131. H. Yang and W. J. van Ooij, Prog. Org. Coat., 50 (2004), 149. 132. B. Chico, J. Simancas, J. M. Vega, N. Granizo, I. Díaz, D. de la Fuente and M. Morcillo, Prog. Org. Coat., 61 (2008), 283. 133. G. Williams and H. N. McMurray, Electrochem. Solid State Lett., 7 (2004), B13. 134. T. Sugawa and K. Gawlik, Mater Lett., 57 (2003), 4282–4290. 135. A. Seth, W. J. van Ooij, P. Puomi, Z. Yin, A. Ashirgade, S. Bafna and C. Shivane, Prog. Org. Coat., 58 (2007), 136. 136. R. G. Buchheit, M. D. Bode and G.E. Stoner, Corrosion, 50 (1994), 205. 137. R. G. Buchheit, S. B. Mamidipally, P. Schmutz and H. Guan, Corrosion, 58 (2002), 3. 138. R. B. Leggat, W. Zhang, R. G. Buchheit and S. R. Taylor, Corrosion, 58 (2002), 322. 139. S. Böhm, H. N. McMurray, S. M. Powell and D. A. Worsley, Materials and Corrosion Werkstoffe und Korrosion, 52 (2001), 896. 140. R. L. Howard, I. M. Zin, J. D. Scantlebury and S. B. Lyon, Prog. Org. Coat., 37 (1997), 83–90. 141. A. Nazarov, D. Thierry, T. Prosek and N. Le Bozec, J. Electrochem. Soc., 152(7) (2005), B220. 142. C. Corfias, N. Pebere and C. Lacabanne, Corros. Sci., 42 (2000), 1337. 143. R. C. MacQueen and R. D. Granata, Prog. Org. Coat., 28 (1996), 97–112. 144. T. Prosek and D. Thierry, Prog. Org. Coat., 49 (2004), 209–217. 145. K. A. Yasakau, M. L. Zheludkevichz and M. G. S. Ferreira, J. Electrochem. Soc., 155(5) (2008), C169–C177. 146. J. Creus, F. Brezault, C. Rebere and M. Gadouleau, Surf. Coat. Technol., 200 (2006), 4636–4645. 147. S. Hayes, P. Yu, T. J. O’Keefe, M. J. O’Keefe and J. O. Stover, J. Electrochem. Soc., 149 (2002), C623–C630. 148. M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, 194. NACE International, Houston, TX, 1974. 149. S. Böhm, R. Greef, H. N. McMurray, S. M. Powell and D. A. Worsley, J. Electrochem. Soc., 147(9) (2000), 3286–3293. 150. F. H. Scholes, C. Soste, A. E. Hughes, S. G. Hardin and P. R. Curtis, Appl. Surf. Sci., 254 (2006), 1770. 151. L. Arurault, P. Monsang, J. Salley and R. S. Bes, Thin Solid Films, 466 (2004), 75. 152. E. B. Li and G. E. Thompson, J. Electrochem. Soc., 146 (1999), 1809. 153. A. M. Cabral, W. Trabelsi, R. Serra, M. F. Montemor, M. L. Zheludkevich and M. G. S. Ferreira, Corros. Sci., 48 (2006), 3740–3758. 154. D. Lau, A. M. Glenn, A. E. Hughes, F. S. Scholes, S. G. Hardin and T. H. Muster, submitted 2008. 155. A. J. Aldykewicz, H. S. Isaacs and A. J. Davenport, J. Electrochem. Soc., 142 (1995), 3342–3350.

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156. A. Braig, Proc. Electrochem. Soc. 97-41(Advances in Corrosion Protection by Organic Coatings), (1998), 18–31. 157. S. B. Saidman, J. Electroanal. Chem., 534(1) (2002), 39. 158. A. Franquet, C. Le Pen, H. Terryn and J. Vereecken, Electrochim. Acta, 48(9) (2003), 1245. 159. V. Branzoi, F. Golgovici and F. Branzoi, Mater. Chem. Phys., 78(1) (2003), 122. 160. S. V. Lamaka, M. L. Zheludkevich, K. A. Yasakau, M. F. Montemor and M. G. S. Ferreira, Electrochim. Acta, 52(25) (2007), 7231. 161. M. L. Zheludkevich, K. A. Yasakau, S. K. Poznyak and M. G. S. Ferreira, Corros. Sci., 47(12) (2005), 3368. 162. B. R. W. Hinton, N. Drubule, A. E. Hughes, M. Forsyth, T, Markley, D. Ho, F. H. Scholes and S. A. Furman, ATB Metall., 45 (2006), 392–401. 163. F. H. Scholes, A. E. Hughes, D. Jamieson, K. Inoue, S. A. Furman, T. H. Muster, S. G. Hardin, D. Lau, T. G. Harvey, P. Corrigan, M. Glenn, P. A. White, J. Mardel and M. Forsyth, Corros. Sci., accepted. 164. T. J. Carney, P. Tsakiropoulos, J. F. Watts and J. E. Castle, !nt. J. Rap. Sol., 5 (1990), 189–217. 165. C. Lea and C. Molinari, J. Mater. Sci., 19 (1984), 2336–2351. 166. A. Roberts, D. Engelberg, Y. Liu, G. E. Thompson and M. Alexander, Surf. Interf. Anal., 33 (2002), 697–703. 167. S. K. Toh, D. G. McCulloch, J. du Plessis, P. J. K. Paterson, A. E. Hughes, D. Jamieson, B. Rout, J. M. Long and A. Stonham, Surf. Rev. Lett., 10 (2003), 365–372. 168. R. K. Viswanadham, T. S. Sun and J. A. S. Green, Corrosion, 36 (1980), 275–278. 169. J. H. Schofield, J. Electron Spectrosc. Relat. Phenom., 8 (1976), 129.

10 Corrosion inhibiting cerium compounds for chromium-free corrosion protective coatings on AA 2024 Michael Schem, Thomas Schmidt, Hinka Caparrotti, Matthias Wittmar and Michael Veith INM — Leibniz Institute for New Materials, Campus D2 2, D-66123 Saarbruecken, Germany [email protected]

10.1

Introduction

AA 2024-T3 is a frequently used aluminium alloy in the aerospace industry. It combines relatively high tensile strength (11 times higher than pure aluminium) with low density. The alloy consists of up to 4.9 wt% copper, 0.9 wt% manganese, 1.8 wt% magnesium, 0.5 wt% silicon, 0.5 wt% iron, and aluminium [1]. The high copper content makes the material very sensitive to corrosion attack since the copper-rich intermetallics act as local cathodes which promote the anodic dissolution of the surrounding aluminium [2]. The state-of-the-art in the corrosion protection of aluminium alloys is the use of chromium(VI)-containing coatings. Due to the toxic and carcinogenic nature of chromium(VI), alternative corrosion protection coatings have to be discovered. One approach to replace chromium(VI) is the use of rare earth chlorides. Mostly these chlorides are used as single inhibitors, but Markley et al. found a synergistic effect when cerium chloride is combined with praseodymium chloride in a ratio that resembles the Ce/Pr ratio in mischmetal [3]. Chambers and colleagues are also examining possible synergistic combinations of inhibitors [4,5]. Several investigations [6–10] have used cerium compounds in conversion coatings for the replacement of hexavalent chromium in a variety of aluminium alloys. The work of Chambers and Taylor [4] identified cerium chloride as the third most efficient inhibitor after sodium chromate and sodium metavanadate after 1 day of exposing AA 2024 wire to an electrolyte containing sodium chloride and the inhibitor to be examined. After 7 days of exposure, cerium chloride is the second best inhibitor after sodium metavanadate. Hamdy and Beccaria examined the influence of pre-treatment on the corrosion protection properties of a cerium chloride conversion coating [11]. If the sample was immersed in the cerium chloride solution without a preceding etching step, then cerium chloride was absorbed non-uniformly. The generation of the cerium layer took several hours due to the need for a three-step treatment, in which each step took at least 1 h. Mansfeld et al. proposed the treatment of aluminium in cerium nitrate solutions [12], in which the substrates are kept for several hours at temperatures close to the boiling point of the solution. The long process duration of these two methods appears to be impractical for industrial use, therefore an alternative way to shorten the 184

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Table 10.1

Solubilities of the examined cerium compounds in water

Salt

Solubility (g/100 g) in water at 25°C

Cerium acetate Cerium acetylacetonate Cerium chloride Cerium nitrate Cerium sulphate

26 [21] 0.17 (at 19°C) [22] 49 [23] 176 [21] 19 (at 0°C) [24]

process times has been demonstrated by Bethencourt et al. [13] and Pardo et al. [14] who electrochemically activated the precipitation of the protective layer. In the present study, the use of cerium compounds in a sol–gel derived hybrid organic– inorganic coating is studied. Apart from the use in conversion coatings, cerium compounds are also used successfully in sol–gel-based coatings to enhance their corrosion protection properties on a variety of metallic substrates [10,15–20]. In a recent paper, Moutarlier et al. [15] have claimed that the use of cerium salts in sol–gel-based corrosion protection is limited due to the high solubility of the cerium ions. In all of the papers cited here, the authors limited themselves to the use of cerium chloride or cerium nitrate. In order to examine the influence of alternative anions, in the present study, cerium sulphate, cerium acetate hydrate (cerium Ac) and cerium acetylacetonate (cerium Acac) have been compared to cerium nitrate and cerium chloride. Additionally, attempts were made to limit the solubility of cerium chloride and cerium nitrate with the addition of acetylacetone to form complexes of the cerium ions. Table 10.1 gives the solubilities of the cerium compounds examined in this study. 10.2 10.2.1

Experimental Synthesis of coating material

An inorganic–organic hybrid sol was synthesised according to Ref. 25. In summary, the sol was synthesised by mixing and hydrolysing three solutions that had been prepared separately. Firstly, tetraethoxysilane (TEOS) and methyltriethoxysilane (MTEOS) and SiO2 nanoparticles, as components of the inorganic network, were mixed and hydrolysed by concentrated hydrochloric acid (HCl). Secondly, 3glycidoxypropyltrimethoxysilane (GPTS) was used for creation of the organic substructure and crosslinking with the inorganic network. To assist hydrolysis, 0.1 M hydrochloric acid was added as a catalyst. At this stage, the cerium compounds were added to the solution. Thirdly, the organic substructure of the hybrid material, consisting of 2,2’-bis-(4-hydroxyphenyl)-propane (BPA), was dissolved in an organic solvent. Finally, after completing the hydrolysis, the three solutions were mixed together. Before coating, the organic crosslinking was commenced by adding 1-methylimidazole to start the organic polymerisation. The cerium compounds used for these experiments were cerium acetylacetonate (cerium Acac), cerium nitrate hexahydrate (cerium nitrate), cerium chloride heptahydrate (cerium chloride), cerium acetate hydrate (cerium Ac) and cerium sulphate. Additionally, cerium nitrate and cerium chloride were dissolved in ethanol and mixed

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with an equimolar amount of acetylacetone (acac). These mixtures were also added to the coating material and were labelled cerium nitrate plus acetylacetonate (cerium nitrate–acac) and cerium chloride plus acetylacetonate (cerium chloride–acac), respectively. These compounds were added to the coatings to achieve cerium concentrations of 2.5 wt% and 4 wt%, respectively, relative to the solid coating material. The salts cerium sulphate, cerium Ac and cerium Acac were not totally soluble in the coating material. Therefore, to ensure a homogeneous distribution and comparable salt amounts in all samples, the mixtures were permanently homogenised and all precipitation was prevented. 10.2.2

Substrates and pre-treatments

Single layer sol–gel coatings were deposited on coupons of aluminium alloy 2024-T3. The aluminium samples were degreased with acetone, cleaned with Metaclean T2001 (Chemie Vertrieb, Hannover, Germany) for 15 min, and then etched in an alkaline cleaner (P3 51 Almeco from Henkel KGaA, Germany) (5 min). Subsequently, the substrates were desmutted with Turco Liquid Smutgo NC (Turco Chemie, Germany) (5 min). The coatings were applied by dip-coating with a coating velocity of 9 mm/s and cured at 120°C for 4 h. 10.2.3

Characterisation of the coated samples

The coating thicknesses were measured by a magnetic induction method with a Permascope (Helmut Fischer GmbH & Co, Sindelfingen, Germany). The corrosion protection performance of the samples was characterised by the neutral salt spray test according to DIN 50021. The backs and edges of the samples were covered with adhesive tape and beeswax, respectively. An artificial scratch was applied to the corrosion protection coating following the procedure used by Van Laar [26] to examine the delamination behaviour. The samples were removed from the test once a week to be photographed with a resolution of 0.07 mm per pixel. These images were then used to assess the corrosion protection by evaluating the appearance after approximately 336 h of salt spray testing. This timespan is important for the evaluation of Cr-free primers according to the MIL-A-8625F norm [27]. If the samples did not show pitting after this time, they were further evaluated by measuring the time in the salt spray test until the first occurrence of corrosion. The long-term barrier performance of the coated samples was tested in a 3.5 wt% sodium chloride solution by electrochemical impedance spectroscopy (EIS) using an Ametek (Princeton Applied Research, TN, USA) model 2273 potentiostat. A threeelectrode electrochemical cell with a platinum sheet and a saturated calomel electrode as the counter and reference electrodes, respectively, was used for the measurements. The sample area exposed to the electrolyte was 6.6 cm2 and the signal amplitude was 10 mV with a frequency range between 5 mHz and 485 kHz. Electrochemical impedance spectra were recorded regularly from each coated sample. Impedance values from the low frequency range can provide information about the self-healing properties of a coating-system [28] and for this reason, the impedance values at 0.01 Hz in particular were plotted against immersion time. The possibility of inhibitor release from a coating and migration to a scratched area is a basic requirement for self-healing properties. For this reason, the leaching of cerium out of the coatings after certain periods of time in ultrapure deionised water

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was determined with an Ultima 2 (Horiba Jobin Yvon GmbH, Munich, Germany) inductively coupled plasma/optical emission spectrometer (ICP/OES) at a detection wavelength of 413.4 nm (limit of detection 0.03 mg/l). Each sample was coated on both sides, cut into four pieces and immersed in 50 ml of doubly distilled water. In order to determine the influence of convection in the extraction water, some of the samples were agitated, while a smaller number were not. To avoid the direct contact of non-agitated specimens within the immersion medium, two edges were coated with beeswax and a gap of approximately 1.5 mm was maintained between them. The agitated samples were not stabilised with beeswax. The agitation was done using a circular motion in a Certomat IS (B. Braun Biotech Int.) at 26°C with a velocity of 180 rpm. The exposed coating area for all samples immersed in the medium was between 50 and 65 cm2. At certain intervals (2.5, 24, 96, 120 h), a 10 ml sample of water was removed from the immersion medium and stored in a glass vial. Subsequently, the samples were analysed for cerium content by ICP/OES. In another set of experiments, samples were collected after immersion for 15, 30, 45 and 90 min. The results of the measurements were corrected to an exposed area of 50 cm² in all cases. Special attention was paid to avoid contamination of the water by particles of the coating material. All ICP measurements were conducted repeatedly to ensure that no pieces of coating which may have become separated from the samples were accounted for as released cerium. 10.3

Results

Neutral salt spray corrosion tests on coupons coated with the system without cerium compounds first showed evidence of corrosion after 170 h of exposure. After 336 h, both corrosion in the general area and sub-surface migration at the artificial scratch were clearly visible (Fig. 10.1). A comparison of the coatings containing cerium Ac and cerium Acac after 336 h of salt spray testing is shown in Fig. 10.2. Comparing cerium Ac and cerium Acac with the coating without cerium additions (Figs. 10.1 and 10.2) shows that pitting in the general area is increased with the cerium additions while creepage at the artificial scratch decreases. In the case of cerium chloride, the samples show only limited pitting in the presence of high amounts of the salt after 336 h of salt spray testing and only minor creepage (Fig. 10.3). In comparison to the cerium Acac-containing samples, the amount of creepage at the artificial scratch is slightly increased, but it is smaller than the creepage of cerium Ac specimens. The addition of acetyl acetone to the solution also results in an increase in creepage at the artificial scratch. The coating containing cerium nitrate showed no pitting corrosion after 336 h of neutral salt spray testing in the case of the lower cerium concentration (Fig. 10.4a). With a cerium concentration of 4 wt%, some isolated corrosion marks were visible in the general area (Fig. 10.4b). The same behaviour was exhibited when using a combination of cerium nitrate and acetylacetone in the coating material. In the case of the low cerium content, no corrosion was observed in the tested area (Fig. 4c). With the higher cerium content, isolated corrosion marks were visible (Fig. 10.4d). The extent of creepage was approximately the same for cerium nitrate sample types, and higher than for pure cerium chloride samples. As shown in Fig. 10.5, the corrosion protection properties of the coating material containing cerium sulphate were very limited. After 336 h of salt spray testing, samples of both concentrations exhibited numerous relatively large pits. The lower

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10.1 Coating material without the addition of cerium compounds after 336 h salt spray test

concentration of cerium seems to result in slightly better performance than the higher one. Comparison of cerium sulphate-containing coatings with cerium Acaccontaining ones shows significant creepage at the artificial scratch after 336 h of salt spray testing with cerium sulphate but very little creepage for Ce Acac samples (Fig. 10.2). The coating thickness was in the range of 8 to 11 μm for all samples except those doped with cerium Ac for which it was 6 μm (Table 10.2). This variance would be expected to have some influence on the coating properties, but lies well within the variance of coatings used industrially. The impedance spectra determined in 3.5% sodium chloride solution showed plateaux in the low frequency region (Figs. 10.6 and 10.7). It is generally accepted that such plateaux can be related to the ohmic resistance of the coating. Therefore, the value of the impedance at low frequencies is used here as a measure of the barrier properties of the coatings [28]. The long-term impedance modulus behaviours in the low frequency region of the 4 wt.% cerium salt filled coating systems, measured via EIS, are shown in Fig. 10.8. The highest impedance modulus of around 7 Mohm cm2 at the beginning of the

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10.2 Cerium Ac (a: 2.5 wt% Ce, b: 4 wt% Ce) and cerium Acac (c: 2.5 wt% Ce, d: 4 wt% Ce) filled coatings after 336 h salt spray test

immersion test (after 1 h and 40 min) was exhibited by the coating containing cerium chloride plus acac. This value decreased during the first 320 h of immersion to approx. 3.2 Mohm cm2 and then remained almost constant. Similar behaviour was displayed by the cerium nitrate-containing coating system but the impedance values were shifted from approx. 2.5 Mohm cm2 to lower values and the plateau was attained after only 200 h of immersion. Cerium nitrate plus acac showed a steep decrease in the impedance modulus during the first 100 h of the experiment, then the impedance increased again to almost

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190

10.3 Images of cerium chloride containing coatings (a: 2.5 wt% Ce as cerium chloride, b: 4 wt% Ce as cerium chloride, c: 2.5 wt% Ce as cerium chloride– acac, d: 4 wt% Ce as cerium chloride–acac) after 336 h salt spray test

the initial value and subsequently showed only a very slow decreasing impedance modulus to 1 Mohm cm2 during a testing period of 2500 h (for purposes of clarity, the timescale in Fig. 10.8 ends at 1200 h). Contrary to the good long-term barrier performance of the cerium salt-doped coating systems mentioned above, the coatings doped with cerium Ac, cerium Acac

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Self-healing properties of new surface treatments

10.4 Images of cerium nitrate (a: 2.5 wt% Ce, b: 4 wt.% Ce) and cerium nitrate– acac (c: 2.5 wt% Ce, d: 4 wt.% Ce) filled coating after 336 h salt spray testing

and cerium sulphate exhibited lower initial impedance modulus values as well as faster degradation of the barrier properties with time. The results of the leaching experiments are presented in Fig. 10.9. Leaching of cerium was detected in the case of the samples doped with cerium Ac, cerium Acac,

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192

10.5 Cerium sulphate-filled coatings after 336 h salt spray testing. (a) 2.5 wt.% Ce, (b) 4 wt% Ce Table 10.2

Measured coating thicknesses

Sample name Unfilled coating Cerium Ac Cerium Acac Cerium chloride Cerium chloride–acac Cerium nitrate Cerium nitrate–acac Cerium sulphate

Cerium concentration (wt%)

Coating thickness (μm)

– 2.5 4 2.5 4 2.5 4 2.5 4 2.5 4 2.5 4 2.5 4

9 6 6 8 8 8 10 9 10 10 11 10 11 9 9

and cerium sulphate. Traces of cerium just above the limit of detection were determined in the case of the higher concentration of cerium chloride and cerium nitrate in the coating. The other coating systems were apparently not able to leach out a detectable amount of their cerium content. In order to determine the influence of the agitation and the resulting convection, Fig. 10.10 compares the detected cerium concentrations of samples that were agitated

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10.6 Impedance spectra (Bode plots) of coatings prepared with 4 wt% cerium chloride, 4 wt% cerium chloride–acac, 4 wt% cerium nitrate, 4 wt% cerium nitrate–acac

10.7 Impedance spectra (Bode plots) of coatings prepared with 4 wt% cerium sulphate, 4 wt% cerium Acac, 4 wt% cerium Ac

Corrosion inhibiting cerium compounds 194

10.8 Impedance modulus |Z| (at 0.01 Hz) against the immersion time in 3.5 wt.% sodium chloride solution for seven different cerium salt-doped coating systems measured by EIS. The concentration of cerium ions in each coating was calculated to be 4 wt%

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10.9 Contents of cerium as a function of immersion time in aqueous media measured by optical emission spectrometry. All samples were agitated. The numbers indicate the measured cerium concentration after 120 h immersion time. The error bars indicate the standard deviation of the ICP/OES measurements.

10.10 Comparison of the cerium concentration in deionised water after different immersion times with and without agitation. The numbers indicate the measured cerium concentration after 120 h immersion time. The data for the agitated samples are the same as in Fig. 10.9

with samples that were stationery during the immersion time. This comparison was carried out for samples containing cerium Acac, and cerium sulphate since these samples showed the highest cerium concentrations and the differences were expected to be greater. In the case of cerium sulphate, a difference was noticeable. The agitated

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196

10.11 Comparison of the cerium content as a function of agitation and sample area

samples appeared to reach the maximum concentration after 2.5 h of immersion time, and the concentration of cerium was not increased by extending the immersion time, while in the case of the non-agitated samples, the cerium concentration continued to increase with time. The detected cerium concentrations after 120 h immersion time were of the same order of magnitude for agitated and non-agitated samples for each initial cerium concentration in the coating material. When comparing the coatings containing cerium Acac, no difference could be identified; the agitated samples showed increases in cerium concentration with immersion time comparable to the non-agitated ones. To verify that the quantity of leached salt was dependent on the area exposed to the pure water, a test was conducted in which only half the surface area was exposed. The results presented in Fig. 10.11 show that the amount of cerium is limited by the sample area exposed to the solution. With an assumed density of the coating material of 1.5 g/cm³, the amount of cerium provided by the coating material could be calculated. The results of this calculation have been related to the maximum concentration of cerium detected by ICP in Table 10.3. The fraction of detected material was in the vicinity of about 50% in all cases. No dominant influence of agitation could be concluded from these data. Table 10.3 Fraction of cerium detected relative to the amount of cerium provided by the coating material

Ce2(SO4)3 4% agitated Ce2(SO4)3 4% Ce2(SO4)3 4% agitated Ce2(SO4)3 4%

Sample area (cm²)

Percentage of cerium detected in eluent

32 26.8 64 54.4

54.4 47.7 45.2 53.3

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10.12 Content of cerium in the diluent after shorter time periods. The numbers indicate the cerium concentration after 90 min of immersion

To investigate the rate at which the cerium concentration developed during leaching, a study was conducted with relatively short immersion times, viz. 5, 15, 45 and 90 min. Cerium Ac, cerium Acac, and cerium sulphate were tested with a reduced leaching time since these inhibitors had leached readily from the coating material in the experiments with longer immersion times (Fig. 10.12). In the time range from 5 to 90 min, in all cases, a progressive increase in cerium concentration was clearly detectable. In the case of cerium sulphate, the final concentration was almost attained after 90 min of leaching. For the other salts, almost 70% of the final concentration was attained within 90 min. 10.4

Discussion

The cerium compounds examined in this study were selected to compare cerium compounds of different solubilities with cerium nitrate, which was proposed for corrosion protection on aluminium by Mansfeld et al. [12]. Cerium chloride and cerium nitrate are soluble in water, while cerium Ac, cerium Acac and cerium sulphate have only a limited solubility in water. According to the data from Table 10.1, the cerium compounds can be arranged in order of increasing solubility as follows: Ce Acac < Ce2(SO4)3 < Ce Ac < CeCl3 < Ce(NO3)3 In order to determine the influence of the chelating effect of acetylacetone further, cerium chloride and cerium nitrate were reacted with acetylacetone. If this reaction had taken place completely, the samples would have shown a similar performance in the salt spray test as the coating filled with cerium Acac. Two major influences of the cerium salts are conceivable: one is an active inhibiting role; the other is further densification of the coating material to increase its barrier properties.

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In the case of cerium compounds with good solubility in the matrix such as cerium chloride and cerium nitrate, the reason for the good corrosion protection can be found in the lack of leaching into the solute. Because of the inability to leach, the coating material is not weakened, and, therefore, the barrier properties remain intact. Due to dispersion on a molecular scale, it could be expected that the pore size in these coatings would be smaller than in the case of matrix with insoluble salts which results in a denser matrix and better barrier properties (see Figs. 10.3 and 10.4). As reported by Bhattamishra and Banerjee [7], in the case of cerium sulphate, it is necessary that a critical concentration of about 500 ppm inhibitor in aqueous 3.5% NaCl solution must be exceeded to facilitate efficient corrosion protection, and further that low amounts of cerium nitrate even increase the corrosion behaviour of aluminium. In the leaching experiments, it turned out that even in the case of cerium sulphate, which exhibited the highest cerium concentration in the leaching experiments, the amount of inhibitor in the diluent was below the limit postulated by Bhattamishra and Banerjee. It is remarkable that the use of cerium minimises the tendency for delamination of the coating material in all cases, while in the case of the unfilled system (Fig. 10.1), a general delamination is observable; only increased creepage can be found at the artificial scratch of most cerium doped samples. On the basis of the salt spray results, the salts can be arranged in order of increasing barrier properties as follows: Ce Acac = Ce2(SO4)3 = Ce Ac < pure matrix << CeCl3 < Ce(NO3)3 and according to the decrease in estimated creepage at the artificial scratch as: Ce2(SO4)3
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cerium species. While cerium sulphate reaches a constant cerium concentration within 90 to 150 min, cerium Acac leaches from the coating material for more than 5 days. This behaviour should be influenced by the limitation of the cerium release by the coating material. The data indicate that only a certain fraction of cerium can be released from the coating material. This might be a result of the way the salts are incorporated into the matrix. Because of the limited solubility of Ce Acac, Ce Ac and cerium sulphate in the matrix, it could be accepted that these coatings form a less dense matrix with larger pores. According to the results, almost 50% of the salt material can be reached through these pores and dissolved by the outer leaching medium. On the other hand, regarding Fig. 10.12, the kinetics of the corrosion processes seem to be slower than the leaching of the salt material. This is responsible for the higher pitting rates of these samples since, after relatively short times, the inhibitor is depleted from the coating and cannot prevent the corrosion processes. With regard to the results of the SST, leaching, and EIS measurements, the best performance of all samples examined was exhibited by the cerium nitrate plus acac coating system. Besides the good performance in the salt spray test, only this coating system showed a significant recovery of the impedance modulus value after an initial decrease and at the same time only a slight decrease in that value with time. Together with the other samples of cerium nitrate and cerium chloride and cerium chloride acac, these materials showed long-term inhibition and stabilisation of the barrier properties. In particular, the increase in the impedance modulus and the nearly constant value might be seen as indicative of self-healing capabilities of these coatings. Other authors have also found that cerium nitrate has good performance as an inhibitor in sol–gel-based coatings [19]. Yasakau et al. found that cerium nitrate did not affect the stability of the sol–gel coating and led to active corrosion protection [29]. The same effective corrosion inhibiting properties of cerium nitrate incorporated into a sol–gel coating was also found by Wang et al. [30]. Kasten et al. investigated cerium nitrate and cerium chloride doped sol–gel coatings and found promising behaviour for these coatings in terms of corrosion protection [31]. Sugama used cerium Ac in a sol–gel matrix to produce cerium oxide in-situ. It was found that the remaining Ac led to less satisfactorily performing coatings [32]. These findings from the literature for cerium Ac, cerium nitrate and cerium chloride are in close correlation to the results discussed in this paper. 10.5

Conclusions

Cerium nitrate and cerium chloride with additions of acetylacetone are, according to the salt spray test results, the most efficient corrosion inhibitors, when used as additives to hybrid organic–inorganic coatings. This was also demonstrated by the results of the EIS measurements. Cerium nitrate and, with reservations, cerium chloride in combination with acetylacetone are, according to the salt spray test performance, possible corrosion protection additives to the hybrid coating system used in this study. The other cerium compounds examined showed only limited corrosion protection abilities according to the results of the electrochemical and salt spray tests. Where a cerium compound was detected in the leaching tests, the performance in the salt spray test was poor. This leads to the conclusion that because of the higher velocity of the salt dissolution in comparison to the kinetics of the corrosion processes, the cerium compound is washed from the coating so that it is not able to exert

Corrosion inhibiting cerium compounds

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its inhibiting function. Additionally, the leaching of the inhibitor from the coating might result in open pores leading to a pathway for corrosive attack. Only coatings with salts soluble in the matrix itself showed long-term preservation of barrier properties in EIS measurements, and in one case a regeneration of a high impedance modulus which could be indicative of self-healing. It can be further concluded that these coatings must have a denser matrix with fewer or smaller pores compared to coatings with less soluble salts that preserve the barrier properties and keep the inhibitor in the place where inhibition should take place. Notwithstanding this study, it is still possible that a coating with good leaching abilities serves well as a primer layer under a topcoat with good barrier properties. In case of damage in the coating, the inhibitors would only leach near the scratch and supply the inhibitor reliably when it is needed. In future, new ways of storing corrosion inhibitors in the coating will be tested, as well as the corrosion protection properties of cerium compounds when the coating material is tested as a primer layer. Acknowledgements The work presented in this paper was funded by the European Community project ‘Multiprotect’ (Contract N° NMP3-CT-2005-011783). References 1. N. A. Belov, D. G. Eskin and A. A. Aksenov, Multicomponent Phase Diagrams: Applications for Commercial Aluminum Alloys, 424; 978-0-08-044537-3. Elsevier, Amsterdam, 2005. 2. K. A. Yasakau, M. L. Zheludkevich, S. V. Lamaka and M. G. S. Ferreira, J. Phys. Chem. B, 110 (2006), 5515–5528. 3. T. A. Markley, A. E. Hughes, T. C. Ang, G. B. Deacon, P. Junk and M. Forsyth, Electrochem. Solid-State Lett., (2007), 10. 4. B. D. Chambers and S. R. Taylor, Corros. Sci., 49 (2007), 1597–1609. 5. B. D. Chambers, S. R. Taylor and M. W. Kendig, Corrosion, 61 (2005), 480–489. 6. B. R. W. Hinton, D. R. Arnott and N. E. Ryan, Mater. Forum, 9 (1986), 162–173. 7. A. K. Bhattamishra and M. K. Banerjee, Z. Metallkd./Materials Research and Advanced Techniques, 84 (1993), 734–736. 8. M. Bethencourt, F. J. Botana, J. J. Calvino, M. Marcos and M. A. Rodríguez-Chacón, Corros. Sci., 40 (1998), 1803–1819. 9. P. Campestrini, H. Terryn, A. Hovestad and J. H. W. d. Wit, Surf. Coat. Technol., 176 (2004), 365–381. 10. H. Schmidt, P. Müller, C. Dittfurth, S. Albayrak and A. Puhl, Sol–Gel Derived Nanocomposite Materials for Corrosion Protection of Aluminium Alloys. 2000 11. A. S. Hamdy and A. M. Beccaria, J. Appl. Electrochem., 35 (2005), 473–478. 12. F. Mansfeld, V. Wang and H. Shih, J. Electrochem. Soc., 138 (1991), L74–L75. 13. M. Bethencourt, F. J. Botana, M. J. Cano and M. Marcos, Appl. Surf. Sci., (2004), 238; 278–281. 14. A. Pardo, M. C. Merino, R. Arrabal, F. Viejo, M. Carboneras and J. A. Munoz, Corros. Sci., 48 (2006), 3035–3048. 15. V. Moutarlier, B. Neveu and M. P. Gigandet, Surf. Coat. Technol., 202 (2008), 2052–2058. 16. A. Pepe, M. Aparicio, A. Duran and S. Cere, J. Sol-Gel Sci. Technol., 39 (2006), 131–138.

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17. A. R. Phani, F. J. Gammel, T. Hack and H. Haefke, Mater. Corros., 56 (2005), 77–82. 18. N. N. Voevodin, N. T. Grebasch, W. S. Soto, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 24–28. 19. M. L. Zheludkevich, R. Serra, K. A. Yasakau, M. G. S. Ferreira, M. F. Montemor and I. M. M. Salvado, Electrochim. Acta, 51 (2005), 208–217. 20. Y. F. Wang, R. M. Wang and Z. C. Guo, Surf. Eng., 24 (2008), 47–51. 21. CRC Handbook of Chemistry and Physics (Internet Version). Taylor and Francis, London, 2007. 22. F. C. Nachod, Z. Elektrochem., 44 (1938), 80–81. 23. M. I. Kadantseva and G. V. Levchenko, Russian Journal of Inorganic Chemistry (Transl. of Zh. Neorg. Khim.), 22 (1977), 1082–1084. 24. Perry’s Chemical Engineer’s Handbook. McGraw-Hill, New York, 1997. 25. H. Schmidt, G. Jonschker and S. Langenfeld, DE 19813709, 1998. 26. J. A. W. van Laar, Dtsch. Farbenzeitschrift, 15 (1961), 56–67 and 104–117. 27. US Department of Defense, 1993. 28. S. V. Lamaka, M. L. Zheludkevich, K. A. Yasakau, R. Serra, S. K. Poznyak and M. G. S. Ferreira, Prog. Org. Coat., 58 (2007), 127–135. 29. K. A. Yasakau, M. L. Zheludkevich, O. V. Karavai and M. G. S. Ferreira, Prog. Org. Coat., 63 (2008), 352–361. 30. H. Wang and R. Akid, Corros. Sci., 49 (2007), 4491–4503. 31. L. S. Kasten, J. T. Grant, N. Grebasch, N. Voevodin, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 11–15. 32. T. Sugama, J. Coat. Technol. Res., 2 (2005), 649–659.

11 Hybrid Ce-containing silica-methacrylate sol–gel coatings for corrosion protection of aluminium alloys Mario Aparicio, Nathaly C. Rosero-Navarro, Yolanda Castro and Alicia Durán Instituto de Cerámica y Vidrio (CSIC), Campus de Cantoblanco, 28049 Madrid, Spain [email protected]

Sergio A. Pellice Instituto de Investigación en Ciencia y Tecnología de Materiales (CONICET). Av. J. B. Justo 4302, B7608FDQ Mar del Plata, Argentina

11.1 11.1.1

Introduction Protection of metal substrates with sol–gel coatings

The sol–gel method is a process in which a precursor in the form of a solution undergoes gelation by evaporation of the solvent and is subsequently cured or sintered to produce a broad range of materials. Industrial applications of the sol–gel process currently focus on protective and functional coatings, utilising the ability of the process to modify substrate surface properties while preserving those of the bulk. This process is attractive because it requires relatively low processing temperatures, gives highly homogeneous products and can be applied to materials with a wide range of compositions and properties. Low viscosity, fluid sols are ideal for preparing films, as long as some limitations are respected. The structure and characteristics of coatings are determined both by the physical properties of the sol (viscosity, surface tension, evaporation rate of the solvent), and by the characteristics of the ‘clusters’ formed from polymerisation of hydrolysed precursors (size, extent of branching or aggregation and fractal dimension of the network). Also, a wide range of coating compositions and properties can be achieved by varying the method for sol preparation (colloidal or hydrolysis– polycondensation), the synthesis parameters (pH, precursors, H2O/precursor ratio, temperature of reaction, sol concentration) and the thermal treatment for consolidation and/or sintering. The sol characteristics and thermal treatments determine a number of the film’s properties, including hardness, critical thickness, wear resistance, chemical behaviour, hydrophobicity and adherence [1–4]. First silica sol–gel coatings were prepared using tetraethylorthosilicate (TEOS) after thermal treatment at 500°C. These pure silica coatings have a critical thickness of 400 nm (measured by profilometry on glass substrates), and they are hard, brittle and dense. The hybrid character is conferred through silicon alkylalkoxides with one or more non-hydrolysable Si–C bonds, or more complex organic groups allowing some polymerisable functionality. The non-hydrolysable Si–C bond allows the structural incorporation of the organic groups. Depending on 202

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the composition of the organic component, it is possible to distinguish between hybrid and hybrid polymer coatings. ‘Hybrid’ coatings, which contain low organic concentrations of highly stable organic groups, are sintered at 350–500°C. ‘Hybrid polymer’ coatings, cured below 200°C, contain higher concentrations of organic groups and usually involve organic and inorganic polymerisations. The number of pathways for sol synthesis can be increased by combining alkoxides and alkylalkoxides, with and/or without polymerisable groups, to obtain the desired properties. Activation of organic polymerisation through epoxy groups or double bonds can be used to produce hybrid networks giving a nanostructured final material consisting of two interpenetrating networks (organic and inorganic) connected by chemical bonds [5]. In the context of corrosion protection, silica layers have the potential to provide a significant improvement because SiO2 has very low oxygen diffusivity and the layers are resistant to corrosion over a wide pH range [6,7]. Furthermore, the incorporation of other oxides (ZrO2, TiO2, Al2O3, B2O3) has increased the range of corrosion applications to alkaline and neutral media for different metal substrates. A complete review of protective sol–gel coatings can be found in the literature [8,9]. Most of the initial shortcomings experienced with pure SiO2 coatings have been overcome by the development of hybrid coatings using methyltriethoxysilane (MTES) in combination with TEOS. It is possible to obtain crack-free coatings up to 2 μm in thickness in one step with a 40TEOS/60MTES molar ratio. The improvement in protection of these hybrid coatings has been attributed to the higher thickness, a lower concentration of defects because of the higher ductility, and the presence of methyl groups that confer higher hydrophobicity, delaying the access of the electrolyte to the substrate surface. While pure and hybrid SiO2 layers are effective as a barrier against oxygen diffusion when the coating is treated at temperatures higher than 350°C, other types of coatings that can be cured or sintered below 200°C are required for metals susceptible to deterioration at higher temperatures. Hybrid polymer coatings meet these requirements [8]. 11.1.2

Lanthanides as environmentally friendly corrosion inhibitors

Chromate conversion coatings (CCCs) provide an excellent protection mechanism against localised corrosion, being the most common option used up to now by industry. CCCs present self-healing behaviour after superficial damage; a scratch or defect appearing in the film can be protected by migration of soluble Cr(VI) species from the coating that precipitate on the surface healing the defect [10]. However, CCCs are extremely dangerous for human health and generate serious problems of environmental contamination. Therefore, it is necessary to find a replacement that supplies self-healing protection against corrosion. A recent trend is the development of sol–gel coatings doped with environmentally friendly inhibitors [11,12]. The aim is to incorporate corrosion inhibitors having the ability to play an active protective role in the case of coating damage. These species would diffuse in the presence of an electrolyte and precipitate on bare metal areas, blocking the corrosion reaction. However, this is not easy to achieve, because the ability of the inhibitor to protect and passivate an active corrosion site on the metal substrate will depend on a balanced combination of properties such as inhibitor efficiency, solubility and diffusivity. Over and above their barrier properties, coatings should behave as reservoirs, slowly releasing the inhibitor

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compound to ensure long lasting protection. This behaviour is intimately related to the precursor form (oxides, ions, nanoparticles) in which the inhibitor has been included in the coating matrix, conditioning the effectiveness of the protective system. Special attention should also be paid to the compatibility between the coating matrix and the inhibitor, because inappropriate entrapment of the inhibitor may lead to performance problems such as blistering caused by osmotic pressure when water migrates within the coating to dissolve the inhibitor, or formation of clusters, which could create new defects [13]. Several routes to introduce a wide variety of inhibitors into sol–gel coatings have been reported in the literature [8]. Lanthanides, especially cerium, fulfil the basic requirements for alternative corrosion inhibitors: the ions form insoluble hydroxides, which enable them to be used as cathodic inhibitors; they have a low toxicity and are relatively abundant in nature. Cerium has a high affinity for oxygen and the bond between cerium and oxygen is unlikely to be broken under the potentials applied. For some aluminium alloys, cerium precipitation from aqueous solutions of cerium salts was observed on cathodic intermetallic compounds and in some instances, the oxide covered the entire specimen surface [14–19]. Hybrid silica coatings prepared with TEOS and MTES doped with Ce(NO3)3 and (NH4)2Ce(NO3)6 on stainless steel and aluminium alloys produce an improvement in corrosion protection with increased immersion time in NaCl solutions. This effect can be explained by the deposition of mixed cerium (III and IV) and chromium hydroxides/oxides on cathodic areas, triggered by an increase in pH. However, thermal treatments at temperatures higher than 350°C are applied to sinter these coatings [20,21]. As a result, these cannot be used to protect metal substrates susceptible to deterioration at lower temperatures. For this reason, most of the silica sol–gel coatings doped with cerium tested on aluminium substrates are epoxy-based hybrid polymer coatings. This composition guarantees good compatibility and bonding with the epoxy-based primer or topcoat typically used with aluminium substrates. Moreover, the availability of silica sol–gel precursors with an epoxy functional group promotes chemical bonding between the sol–gel coating and the substrate. Some studies have incorporated cerium nitrate into the epoxy-based sol–gel coatings prepared by combining glycidoxypropyl trimethoxysilane (GPTMS) and different silicon alkoxides or zirconium(IV) tetrapropoxide (TPOZ) [22–26]. To evaluate the selfhealing properties of cerium-doped coatings on aluminium, an artificial defect was produced on the film with a diamond tipped scribe. Results obtained with epoxy-based sol–gel coatings doped with cerium acetate indicated that the inhibitor interacts with the sol structure, creating particles with cracks that can act as initiation sites for localised corrosion, as a result of which the performance deteriorates with immersion time. Thus, it looks necessary to adapt the sol–gel chemistry to include inhibitors for an improved corrosion protection system. Cerium incorporated in ionic form could diffuse too quickly, reducing the protection with time. On the other hand, stabilising the inhibitor within the coating matrix through chemical bonds could limit its ability to diffuse to active reaction sites. In situ formation of ZrO2 nanoparticles within hybrid coatings using TPOZ was found to enhance the corrosion protection [27]. The presence of these amorphous nanoparticles is believed to have a pore blocking effect that improves the barrier properties of the coating.

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11.2

Design of silica-methacrylate hybrid polymer sol–gel materials

The alternative we proposed for substituting chromate conversion coatings (CCCs) for application on metal substrates treatable up to 150–200°C is based on the development of thick hybrid polymer sol–gel coatings. These coatings have to combine barrier properties to delay the penetration of corrosion agents and inhibition properties to hinder the corrosion process because of the presence of a pore, crack or scratch in the coating. The inhibition action through a self-healing or self-sealing process has several steps: diffusion of the inhibitors (cerium ions in this case) across the coating towards the corrosion sites, precipitation of insoluble products and deposition covering the corrosion area. The efficiency of the cerium ion diffusion depends on the hybrid structure porosity: percentage, size and connectivity. A highly cross-linked and dense structure could provide an excellent barrier against corrosion reactives, such as oxygen, water, Cl–, etc., but a very slow cerium ion diffusion making ineffective the inhibition action. On the other hand, an open hybrid structure would facilitate Ce diffusion, but the barrier function would surely diminish permitting corrosion to proceed quickly, the inhibitor being insufficient to stop the corrosion process. The design of the hybrid polymer structure of the coating for a specific metal substrate and corrosive conditions is crucial to combine adequately both properties, barrier and inhibition, so as to obtain an efficient corrosion protection system. As pointed out above, most of the hybrid polymer sol–gel coatings in the literature are based on silica-epoxy structures because of their highly cross-linked structure and chemical compatibility with the epoxy-based primer and topcoat used in most applications of Al alloys. However, a sol–gel coating based on silica-methacrylate offers the opportunity to design a structure with an adequate level of cross-linking to optimise cerium diffusion to provide long-term corrosion protection. The control of the organic polymerisation, through double bonds, together with inorganic condensation makes possible the design of the concentration of chemical bonds in the structure. The types of precursors used to prepare these hybrid structures are: (1) silicon alkoxides to build the inorganic structure through the hydrolysis and condensation of alkoxide groups; (2) monomers of methacrylate groups to control the organic/inorganic ratio; (3) modified silicon alkoxides with methacrylate groups to bond both components, organic and inorganic, to build a nanostructured material with two interpenetrating networks joined by chemical bonds. These materials belong to the Hybrid class II following the classification of Gómez-Romero and Sanchez [5]. The compensation of the polymerisation (organic and inorganic) kinetics is very important to avoid the formation of a composite material with large organic and inorganic domains and a low connectivity between them. The control of the processing parameters: temperature, pH, concentration and water molar ratio is the key to produce these nanostructured materials. Solutions were synthesised from TEOS, 2-hydroxyethylmethacrylate (HEMA) and 3-methacryloxypropyltrimethoxysilane (MPS), the last one as the coupling agent between the organic and inorganic networks thanks to a double bond and three methoxy groups, respectively (Fig. 11.1). The free-radical copolymerisation of the C=C groups of MPS and those of the co-monomer was carried out by using a suitable initiator (2,2´-azobis(isobutyronitrile), AIBN) [28]. The organic polymerisation and inorganic condensation lead to a sol constituted by microgels of sizes

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11.1 Chemical structures of TEOS, HEMA and MPS

between 100 and 300 nm in an alcoholic medium. The concentration of the sol is the key parameter to control the degree of C=C polymerisation to design a more or less permeable structure, and also to define the stability of the solutions. As an example, a schematic view of the chemical structure of a sol with a concentration of Si of 19 g L–1, an inorganic condensation of 2.86 and a complete organic polymerisation is shown in Fig. 11.2. The initial viscosity of this sol is 2 mPa.s and with a gelation time of around 180 h at 25°C. The sol has excellent properties to produce coatings: good wettability, homogeneity, adequate viscosity and high

11.2 Schematic view of the chemical structure of a 60TEOS-10MPS30HEMA sol with a concentration of Si of 19 g L–1, an inorganic condensation of 2.86 and a complete organic polymerisation

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stability (more than 6 months) at 5°C. The structure also allows the incorporation of corrosion inhibitors, such as cerium ions, due to the presence of hydrophilic groups such as OH groups from HEMA and silanols. 11.3 Silica-methacrylate hybrid polymer sol–gel coatings with self-healing properties for the corrosion protection of aluminium alloys 11.3.1

TEOS-MPS-HEMA coatings

As has been pointed out, TEOS-MPS-HEMA sols have adequate properties to prepare coatings on aluminium alloys. Figure 11.3 shows a SEM photograph of an AA2024 alloy protected by a three-layer 60TEOS-10MPS-30HEMA (molar composition) coating prepared by dip-coating. The coating is homogeneous, crack-free and well adhered to the substrate, with a thickness around 6.5 μm. These coatings provide a small barrier functionality at initial immersion times in NaCl solutions. The low degree of cross-linking in the structure and the highly hydrophilic nature because of the presence of hydroxyl groups from HEMA lead to quick electrolyte permeation and, consequently, to the deterioration of the coating [29]. The incorporation of cerium ions in the sols, through the dissolution of cerium salts, such as cerium chloride (CeCl3·7H2O) and cerium nitrate (Ce(NO3)3·6H2O), gives important modifications in the coating response. On the one hand, the increase in defect concentration due to the disruption of the structure reduces the barrier functionality of the coatings. On the other, although cerium ions provide a self-healing effect precipitating as yellowish oxide-hydroxide, the amount of cerium is not sufficient for the numerous corrosion points. The mechanism

11.3 SEM photograph of an AA2024 alloy coated with a three-layer 60TEOS-10MPS-30HEMA coating

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proposed by the self-healing or self-sealing ability provided by cerium ions is based on the tendency for oxide formation, as the values of free energy (ΔGf) are –1024.6 kJ mol–1 for CeO2 and –1706.2 kJ mol–1 for Ce2O3. The oxidation state of cerium ions in the coatings is mainly the same as in the salt used in the synthesis, although a mixture of Ce3+ and Ce4+ is always present [20,21]. Both cerium ions contained in the coating present a high reactivity with oxygen, being the driving force for cerium migration through the coating to react with OH– groups produced in the cathodes and to form cerium oxides/hydroxides that block the subsequent entrance of oxygen to the reactive site: Ce3+ + 3OH 2Ce(OH )3

11.3.2

Æ ¨

Æ ¨

Ce2O3

Ce 4 +

+ 2OH -

Ce(OH )2 2 +

+ 2OH -

Æ ¨ Æ ¨

Ce(OH )3 + 3H 2O Ce(OH )2 2 + CeO2

+ 2 H 2O

TEOS-MPS-SiO2 coatings

Based on the results in the previous section, a new design of hybrid structure is required to improve the cross-linking and density of the coatings to prevent rapid electrolyte permeation. The new design is supported by two modifications [30]: 1. Removal of HEMA to eliminate its C–OH groups that provide high hydrophilicity but very low bonding capacity. 2. Incorporation of commercial silica nanoparticles (Levasil 200A, Bayer, particle size 20 nm) with surface Si–OH groups to increase the density of the coatings. Figure 11.4 presents Bode plots for AA2024 protected with a 42TEOS-23MPS35SiO2 (molar composition) coating (3.7 μm) after different immersion times in 3.5 wt.% NaCl solution compared with the bare alloy. The phase angle curve of the bare alloy has two time constants at 30 and 0.015 Hz assigned to the intermediate aluminium oxide layer and the electron charge transfer process from corrosion, respectively [27]. The incorporation of the sol–gel coating gives a new time constant at higher frequencies (above 105 Hz). The sol–gel coating leads to an increase in the impedance modulus at 0.01 Hz by two orders of magnitude as a result of the additional barrier functionality. However, Bode plots of coated samples after only 1 h of immersion showed signs of corrosion activity by the presence of a time constant at 0.02 Hz. Although the modification of the structure produces an improvement in the barrier functionality compared with TEOS-MPS-HEMA coatings, a porous structure remains in this new sol–gel coating explaining the presence of corrosion after 1 h of immersion in NaCl solution. On the other hand, a relatively open structure like this can be adequate for the objective of combining barrier properties and self-healing effect when an inhibitor such as cerium is incorporated in the coating. The increase in the immersion time produced a slow deterioration of the corrosion protection system. The total impedance at 0.01 Hz decreases slightly with time as a first sign of degradation. In addition, the plateau observed in the impedance plot between 1 and 100 Hz, associated with the contribution of the resistances assigned to the NaCl

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11.4 Bode plots for AA2024 protected with 42TEOS-23MPS-35SiO2 coatings after different immersion times in 3.5 wt.% NaCl solution compared with the bare alloy

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solution, sol–gel coating and intermediate layer, decreases with immersion time, also indicating coating degradation. The reduction of phase angle of the two higher frequency time constants with immersion time indicates a less capacitive response due to the permeation of the solution through the pores of the sol–gel coating and intermediate layer. The incorporation of cerium ions from a salt, Ce(NO3)3·6H2O, reduces the barrier functionality, but generates an inhibition process. Figure 11.5 presents Bode plots for AA2024 protected with a 40TEOS-22MPS-33SiO2-5Ce (molar composition) coating (4.3 μm) after different immersion times in 3.5 wt.% NaCl solution. The values of impedance at 0.01 Hz are almost two orders of magnitude lower than those obtained with sol–gel coatings without cerium and close to values for the bare substrate. This change is an indication of a coating with a higher content of defects [31]. Although it is not as evident as in the case of coatings without cerium, the phase angle curves maintain the three time constants already observed and assigned to the sol–gel coating, intermediate aluminium oxide layer and electron charge transfer process. The presence of diffusion processes through pores is exhibited by a phase angle shift of 45° in the phase plot and a decreasing slope in the impedance modulus plot at lower frequencies. The reduction of the impedance at around 103 Hz is also a sign of the deterioration of the barrier functionality provided by the sol–gel coating and intermediate layer. Although the incorporation of cerium ions has induced a partial degradation of the coating as a barrier, an inhibition action has also been provided. A first sign of this inhibition mechanism is the increase in the impedance modulus at 0.01 Hz with immersion time, contrary to that observed in coatings without cerium. As resistances assigned to the sol–gel coating and intermediate layer decrease with immersion time, the rise of the impedance at 0.01 Hz can only be associated with a significant increase in the resistance describing the corrosion of the metal substrate (Rct). This resistance is inversely proportional to the corrosion rate, and should be normalised with respect to the electrode area. Rct can be defined as: Rct = Rct0/Acorr Rct0 is the Rct value of the metal substrate without coating, and Acorr is the substrate area affected by corrosion [32]. On the other hand, the increase in the phase angle from 30° for 1 h to almost 80° after 215 h of immersion indicates a more capacitive behaviour. Both factors (increase in Rct and phase angle) indicate the presence of deposits or precipitates well adhered and with enough density to reduce the area affected by corrosion. Figure 11.6 shows photographs of the tested areas of coated samples after the EIS tests. The alloy protected with the 42TEOS-23MPS-35SiO2 coating (Fig. 11.6a) presents a transparent and almost intact film with only scarce isolated pitting. This coating provides a good barrier, although the initiation of pitting cannot be stopped due to the absence of a corrosion inhibitor. Figure 11.6b shows a photograph of the sample protected with the 40TEOS-22MPS33SiO2-5Ce coating after the EIS-time test. Well adhered yellowish precipitates can be observed all over the test area. Although the initiation of a wide corrosion process is evident, macroscopic pitting, such as that observed on samples protected with coatings without cerium, has not been detected. This behaviour can be attributed to the self-healing effect provided by the cerium.

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11.5 Bode plots for AA2024 protected with 40TEOS-22MPS-33 SiO2-5Ce coatings after different immersion times in 3.5 wt.% NaCl solution

The improvement in the barrier functionality, compared with TEOS-MPS-HEMA coatings, and the confirmation of the inhibition properties of cerium ions have been shown. However, the increase in defect concentration based on the disruption of the hybrid structure makes necessary the re-design of the coating network and the model of a single layer.

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11.6 Photographs of the tested areas of samples with coatings after the EIS-immersion time analysis: (a) 42TEOS-23MPS-35SiO2 coating and (b) 40TEOS-22MPS-33SiO2-5Ce coating

11.3.3

TEOS-MPS-EGDMA-SiO2 multilayer coatings

The model of a protective coating with a single layer that combines barrier functionality with the inhibition action of cerium ions has been demonstrated to be an extremely difficult objective. For this reason, a new design was based on the synthesis of a multilayer coating where some layers do not contain cerium ions and have to perform only as a barrier against diffusion of corrosive agents, while other layers doped with cerium ions should provide a self-healing effect when the barrier is damaged. A new monomer, ethylene glycol dimethacrylate (EGDMA), is incorporated into the structure to increase the network cross-linking through the polymerisation of its two double bonds (Fig. 11.7). Figure 11.8 compares the polarisation curves after 1 h of immersion in 0.05 M NaCl of the bare AA2024, the alloy protected with a 42TEOS-23MPS-35SiO2 monolayer coating, and the same alloy protected with a tri-layer coating formed by an inner and outer 42.5TEOS-17MPS-8.5EGDMA-32SiO2 coating and an intermediate 40.5TEOS-16.2MPS-8.1EGDMA-30.6SiO2-4.6Ce coating. The bare alloy presents active dissolution of the metal substrate in the anodic branch,

11.7 Chemical structure of ethylene glycol dimethacrylate (EGDMA)

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11.8 Polarisation curves after 1 h of immersion of AA2024 alloys in 0.05 N NaCl: (A) bare, (B) alloy protected with a 42TEOS-23MPS-35SiO2 monolayer coating and, (C) alloy protected with a tri-layer coating formed by an inner and outer 42.5TEOS-17MPS-8.5EGDMA-32SiO2 coating, and an intermediate 40.5TEOS-16.2MPS-8.1EGDMA-30.6 SiO2-4.6Ce coating

without signs of passivation. The coatings provide a good barrier against corrosion, both showing an icorr value decreasing by almost three orders of magnitude compared to the bare alloy. The difference in the coating thickness, 3.7 μm for the 42TEOS-23MPS-35SiO2 coating and 11 μm for the tri-layer coating, can be distinguished in the passive region shown in the anodic branch. The former presents a passive zone of 500 mV compared with the 1200 mV of the tri-layer coating. The EIS analysis clarifies the role and performance of the different coatings. The impedance study (EIS) as a function of the immersion time in 0.05 M NaCl of a bi-layer coating (8 μm) prepared with 42.5TEOS-17MPS-8.5EGDMA32SiO2 sol is displayed in Fig. 11.9. The curve for 1 h of immersion also reflects the barrier properties of these coatings, because only two time constants can be observed at around 104 and below 10–2 Hz, assigned to the sol–gel coating and the alumina layer, respectively. The first signs of corrosion activity can be seen after 17 h of immersion where three time constants appear. The increase in immersion time gives a continuous slow deterioration of the protection system with a reduction of the impedance over the entire frequency range, mainly the impedance associated with the sol–gel coating. In spite of the good barrier properties of these coatings, the absence of an inhibitor allows the corrosion process initiated in any defect of the coating to proceed. Figure 11.10 shows Bode plots after different immersion times in 0.05 M NaCl solution of AA2024 protected with a tri-layer coating (11 μm) formed by an inner and outer layer prepared using a 42.5TEOS-17MPS-8.5EGDMA-32SiO2 sol, and an intermediate layer prepared using a 40.5TEOS-16.2MPS-8.1EGDMA30.6SiO2-4.6Ce sol. The results show a very good performance against corrosion. First signs of corrosion activity only appear after 1870 h of immersion with a time constant centred around 0.3 Hz. The stability of the impedance value at 0.01 Hz

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11.9 Bode plots of AA2024 protected with a bi-layer coating prepared with the 42.5TEOS-17MPS-8.5EGDMA-32 SiO2 sol after different immersion times in 0.05 N NaCl solution

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11.10 Bode plots of AA2024 protected with a tri-layer coating formed by an inner and outer layer prepared using a 42.5TEOS-17MPS-8.5EGDMA-32 SiO2 sol, and an intermediate layer prepared using a 40.5TEOS-16.2MPS8.1EGDMA-30.6 SiO2-4.6Ce sol after different immersion times in 0.05 N NaCl solution

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after 400 h of immersion indicated a significant difference compared with the previous coating without cerium. This behaviour could be related to the healing of the first corrosion sites. The presence of a new time constant at 80 Hz after 3000 h of immersion could be related to the formation of cerium-based precipitates. This finding is under study. Figure 11.11 displays several photographs of the test area of protected AA2024 alloy with both types of coating after different immersion times in 0.05 M NaCl solution. The images clearly show the influence of the incorporation of cerium ions in the coatings. Figures 11.11A and 11.11B present photographs of the alloy protected by the bi-layer 42.5TEOS-17MPS8.5EGDMA-32SiO2 coating, showing the continuous degradation of the coating with immersion time. Although this type of coating has good barrier properties, the slow penetration of water through the defects cannot be stopped because of the absence of inhibitors. On the other hand, the coatings doped with cerium present a very different behaviour. Figures 11.11C and 11.11D show photographs of the alloy protected by the tri-layer coating formed by an inner and outer 42.5TEOS-17MPS8.5EGDMA-32SiO2 layer, and a 40.5TEOS-16.2MPS8.1EGDMA-30.6SiO2-4.6Ce intermediate layer. In this case, there is not evidence of aggressive corrosion, and the nature of the coating surface is unchanged with immersion time. Only a large pit appears in the surface, and it seems to be stabilised by the effect of the cerium ions. The SEM study of this sample after 3130 h of immersion (Fig. 11.12) shows a cracked coating with the presence of precipitates. EDS analysis shows a very significant difference in cerium content between the precipitate and the coating, taking into account that the nominal Si/Ce molar ratio is 95:5. This is further evidence of cerium diffusion and precipitation, probably in cathode regions, delaying the corrosion process.

11.11 Photographs of the test area of protected AA2024 alloy with both types of coating after different immersion times in 0.05 N NaCl solution: (A) bi-layer coating without cerium after 1870 h; (B) bi-layer coating without cerium after 2613 h; (C) tri-layer coating with cerium after 1870 h; (D) tri-layer coating with cerium after 3130 h

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11.12 SEM micrograph of the tri-layer coating with cerium after 3130 h showing some cracks and a precipitate. The EDS results of the Si/Ce molar ratio are included for two locations

These examples of Ce-containing silica-methacrylate sol–gel coatings on Al2024 alloy provide different evidence of the appearance of self-healing or self-sealing effects related to the presence of Ce ions. A major objective of further research now in progress is to increase the Ce content in the film with the aim of obtaining coatings acting as reservoirs of inhibitors for a more prolonged corrosion protection effect. 11.4

Conclusions

This study describes the evolution of the structural design of cerium-doped silica-methacrylate hybrid polymer sol–gel coatings to protect aluminium alloys combining barrier functionality and inhibition properties. The coatings prepared using tetraethylorthosilicate (TEOS), 2-hydroxyethylmethacrylate (HEMA) and 3-methacryloxypropyltrimethoxysilane (MPS) provide a minor barrier functionality, although the incorporation of cerium ions gives a self-healing effect precipitating as yellowish oxide-hydroxide. Another sign of this inhibition mechanism is the increase in the impedance modulus at 0.01 Hz with immersion time, contrary to what is observed in coatings without cerium. Several modifications such as removal of HEMA, and incorporation of commercial silica nanoparticles and ethylene glycol dimethacrylate (EGDMA), produce an increase in crosslinking and density with a significant improvement in barrier functionality. However, the only way to combine barrier functionality and self-healing effect in this system is to develop a multilayer coating where each layer has a specific role. The results present a very good behaviour against corrosion with the barrier showing the first signs of corrosion activity after 1870 h. On the other hand, the stability of the impedance values at 0.01 Hz after 400 h of immersion indicated a significant difference compared with the previous coating without cerium. This behaviour could be related to the healing of the first corrosion sites produced.

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After the EIS-time test, the alloy protected with the coating without cerium showed a general degradation, while the coatings doped with cerium did not show evidence of aggressive corrosion. In this case, only a large pit appeared in the surface, and it seemed to be stabilised by the effect of the cerium ions. The EDS analysis showed a very significant difference in cerium content between the precipitate and the coating, evidencing the cerium diffusion and precipitation, probably in cathode regions, delaying the corrosion process. References 1. J. D. Mackenzie, J. Non-Cryst. Solids, 48 (1982), 1–10. 2. C. J. Brinker and G. W. Scherer, Sol–Gel Science: The Physics and Chemistry of Sol–Gel Processing. Academic Press, New York, 1990. 3. J. Zarzycki, J. Sol–Gel Sci. Technol., 8 (1997), 17–22. 4. D. R. Uhlman and G. Teowee, J. Sol–Gel Sci. Technol., 13 (1998), 153–162. 5. P. Gómez-Romero and C. Sanchez, Functional Hybrid Materials. Wiley-VCH, Germany, 2004. 6. O. de Sanctis, L. Gómez, A. Marajofsky, C. Parodi, N. Pellegri and A. Durán, J. NonCryst. Solids, 121 (1990), 315–318. 7. O. de Sanctis, N. Pellegri and A. Durán, Surf. Coat. Technol., 70 (1995), 251–255. 8. A. Durán, Y. Castro, M. Aparicio, A. Conde and J. J. de Damborenea. Int. Mater. Rev., 52 (2007), 175–192. 9. A. Durán, Y. Castro, A. Conde and J. J. de Damborenea, Chapter 19: ‘Sol-gel protective coatings for metals’. Volume III: ‘Applications of sol-gel technology’, Handbook of Sol-Gel Science and Technology, Processing, Characterization and Applications, ed. S. Sakka. Kluwer Academic Publishers, Dordrecht, 2005. 10. J. Zhao, G. Frankel and R. L. McCreery, J. Electrochem. Soc., 145 (1998), 2258–2264. 11. R. F. Brady, Jr and R. W. Drisko, Kirk–Othmer Encyclopedia of Chemical Technology, Vol. 6, 4th edn, 748. John Wiley and Sons, Inc, New York, 1993. 12. Y. J. Du, M. Damron, G. Tang, H. Zheng, C. J. Chu and J. H. Osborne, Prog. Org. Coat., 41 (2001), 226–232. 13. J. H. Osborne, K. Y. Blohowiak, S. R. Taylor, C. Hunter, G. Bierwagon, B. Carlson, D. Bernard and M. S. Donley, Prog. Org. Coat., 41 (2001), 217–225. 14. M. A. Arenas, M. Bethencourt, F. J. Botana, J. de Damborenea and M. Marcos, Corros. Sci., 43 (2001), 157–170. 15. M. Bethencourt, F. J. Botana, J. J. Calvino, M. Marcos and M. A. Rodriguez Chacon, Corros. Sci., 40 (1998), 1803–1819. 16. M. A. Arenas, A. Conde and J. J. de Damborenea, Corros. Sci., 44 (2002), 511–520. 17. B. R. W. Hinton and L. Wilson, Corros. Sci., 29 (1989), 967–975. 18. R. Di Magio, L. Fedrizzi, S. Rossi and P. Scardi, Thin Solid Films, 286 (1996), 127–135. 19. P. P. Trzaskoma-Paulette and A. Nazeri, J. Electrochem. Soc., 144 (1997), 1307–1310. 20. A. Pepe, M. Aparicio, A. Durán and S. Ceré, J. Sol–Gel Sci. Technol., 39 (2006), 131–138. 21. A. Pepe, M. Aparicio, S. Cere and A. Duran, J. Non-Cryst. Solids, 348 (2004), 162–171. 22. W. Trabelsi, E. Triki, L. Dhouibi, M. G. S. Ferreira, M. L. Zheludkevich and M. F. Montemor, Surf. Coat. Technol., 200 (2006), 4240–4250. 23. W. Trabelsi, P. Cecilio, M. G. S. Ferreira and M. F. Montemor, Prog. Org. Coat., 54 (2005), 276–284. 24. N. N. Voevodin, N. T. Grebasch, W. S. Soto, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 24–28. 25. L. S. Kasten, J. T. Grant, N. Grebasch, N. Voevodin, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 11–15.

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26. M. L. Zheludkevich, R. Serra, M. F. Montemor, K. A. Yasakau, I. M. Miranda Salvado and M. G. S. Ferreira, Electrochim. Acta, 51 (2005), 208–217. 27. M. L. Zheludkevich, R. Serra, M. F. Montemor, I. M. Miranda Salvado and M. G. S. Ferreira, Surf. Coat. Technol., 200 (2006), 3084–3094. 28. S. A. Pellice, R. J. J. Williams, I. Sobrados, J. Sanz, Y. Castro, M. Aparicio and A. Durán, J. Mater. Chem., 19 (2006), 3318–3325. 29. D. A. López, N. C. Rosero-Navarro, J. Ballarre, A. Durán, M. Aparicio and S. Ceré, Surf. Coat. Technol. (in press). 30. N. C. Rosero-Navarro, S. A. Pellice, A. Durán and M. Aparicio, Corr. Sci. (in press). 31. M. L. Zheludkevich, D. G. Shchukin, K. A. Yasakau, H. Möhwald and M. G. S. Ferreira, Chem. Mater., 19 (2007), 402-411. 32. A. Amirudin and D. Thierry, Prog. Org. Coat., 26 (1995), 1–28.

12 Organosilicon plasma polymer coatings containing Ce-based nanoparticles: characterisation of anti-corrosion properties Doriane Del Frari, Jérôme Bour, Julien Bardon, Olivier Buchheit, Claire Arnoult and David Ruch Centre de Recherche Public Henri Tudor (CRPHT), Laboratoire de Technologies Industrielles, Rue de Luxembourg 66, L-4002 Esch-sur-Alzette, Luxembourg, Belgium [email protected]

12.1

Introduction

For many decades, chromate compounds have been successfully used as anticorrosive inhibitors in the surface treatment of aluminium and others alloys. The use of chromates is however restricted worldwide, as they are considered highly toxic and carcinogenic [1]. This has stimulated research aimed at the development of effective and environmentally acceptable alternatives to chromates. To obtain the beneficial properties of chromatation, two approaches can be utilised: passive [2] and active [3,4] corrosion protection. Passive protection is normally provided by a barrier film that prevents contact between the corrosive species and the metal surface and therefore hinders a corrosion process. However, when a defect is formed in the barrier layer, the coating cannot stop corrosion in that place. The second approach is active corrosion protection, which employs inhibitive species that can decrease corrosion activity. An important point is that both strategies must be used together to protect the metallic substrate adequately. Many alternative processes have been examined: among the possible candidates for the environmentally friendly protection of alloys are the organosilicon-based treatments [5], particularly for the protection of galvanised steel [5,6]. Dry processes such as plasma discharge deposition, could be preferred because they are solvent-free and, therefore, more environmentally friendly. Organosilicon layers such as plasma polymerised hexamethyldisiloxane (ppHMDSO) coatings can be deposited on metal parts (e.g. galvanised steel) by Dielectric Barrier Discharge (DBD) atmospheric plasma processes [7]. According to the literature, polymerised organosilicon-based coatings are known to be effective treatments to achieve physical barriers layers [8]. The major drawback of the siloxane coatings is their inert character regarding the corrosion processes. By themselves, ppHMDSO coatings do not provide any ‘active’ protection when barrier coatings fail and aggressive species reach the metallic surface and initiate corrosion activity. Recent research efforts have been focused on the modification of the bulk properties of organosilicon coatings by adding ‘active’ anti-corrosion species to improve corrosion resistance further, and/or to introduce self-healing ability in the coating [9–13]. It has been reported that the specific incorporation of a small amount (1–5%) 220

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of nanoparticles in organic polymers leads to improvements in the barrier properties of organic coatings. Van Ooij and colleagues [14] have shown that nanoparticle-filled silane films could be used to replace the toxic chromate-based treatment for aluminium alloys. Silane films can be thickened and strengthened by incorporating silica nanoparticles into the film. The authors suggest that the presence of silica suppresses the cathodic reaction (oxygen reduction). The reaction between hydroxyl ions and silica leads to the formation of SiO3– ions, which in a second step, react with Al3+ ions to form a passive silica film. Some studies have examined the possibility of combining plasma polymerisation techniques and nanocomposite elaboration methods. Nastase and colleagues [15–17] worked on nanocomposites with nanoparticles and nanotube fillers incorporated into a polymer matrix by plasma polymerisation. This process allows direct and easy deposition of the nanocomposite coating on the target surface. They demonstrated that this technique could be competitive with chemical synthesis and sol–gel processes. The development of anticorrosion coatings with properties similar to those of hexavalent chromium coatings which provide both self-healing and physical barrier effects is challenging. It is known that cerium-based coatings can also provide corrosion protection of metals thanks to their active properties. Indeed, cerium oxides and cerium hydroxides are reported to be cathodic inhibitors and have also been proposed as effective species for the protection of metals from corrosion: cerium possesses similar behaviour to that of chromium [18–20]. Hamdy [21] studied the effect of cerium treatment on corrosion behaviour. A silicate/cerate composite treatment has been tested and was found to improve the corrosion resistance due to the formation of an oxide-thickening layer that acts as a barrier to oxygen diffusion to the metal surface. Cerium is incorporated into the pores of the oxide film and is concentrated at the metal/film interface, leading to improved corrosion resistance. Moreover, it seems that a pitting auto-repair process takes place when high amounts of cerium and silicon are present. Additional improvement in corrosion protection was also observed in the case of silane pretreatments when cerium nitrate was introduced into the coating [15]. Therefore, innovative layers composed of siloxane/cerium and deposited by atmospheric pressure plasma have been tested. In this study, HMDSO was atomised and introduced as a precursor in an atmospheric pressure DBD plasma. The hybrid coating was obtained by mixing liquid precursor and nanoparticles (HMDSO and nanoAlCeO3) before atomisation. The properties of these different coatings were studied by scanning electron microscopy (SEM) and interferometry measurements. Their corrosion resistance has been determined electrochemically and their selfhealing properties have been demonstrated by a combination of electrochemical impedance spectroscopy (EIS) and a nanoscratch method. 12.2 12.2.1

Experimental Substrates and solutions

The samples of hot dip galvanised cold rolled steel (ALUZINC) to be treated were provided by Arcelor Mittal Dudelange (Luxembourg) and the galvanisation coating contained 55 wt% Al, 43.3 wt% Zn and 1.6 wt% Si. They were degreased using ethanol. The chromate coating on the reference sample contained about 20 mg/m2 of hexavalent chromium.

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SEM analysis of the AlCeO3 nanoparticles (Aldrich, 99%) showed them to have diameters varying from a few nanometres (<25 nm) to 300 nm. They were added to a hexamethyldisiloxane (HMDSO > 98% grade) solution, purchased from Aldrich. The concentration of nanoparticles was calculated in weight percent. The solutions were homogenised by ultrasonic stirring at room temperature for 30 min. 12.2.2

Plasma deposition process

Coating experiments were performed in a semi-dynamic dielectric barrier discharge (DBD) reactor. The DBD discharge was switched on between an earthed bottom aluminium plate and two high voltage aluminium top plates protected by a 3.25 mm thick glass plate, as shown in Fig. 12.1. The gap between the earthed electrode and the dielectric plate was set to 2 mm. The liquid precursor solution was atomised at a constant pressure of 2×105 Pa and injected into the carrier gas flow before entering the plasma zone. The carrier gas composition, which was set by mass flow controllers, was a N2/O2 (97:3, vol.%) mixture with a total gas flow of 20 standard litres per minute. The gas mixture containing the precursor aerosol was injected into the plasma through a slit between the two top electrodes. The experiments were carried out at atmospheric pressure and ambient temperature. During the deposition experiments, the top electrode block moved back and forth over the sample at a constant speed (4 m/min) with a 380 mm displacement range. A treatment with 12 passes took 68.4 s. It has been demonstrated that the coating properties can be improved when an additional plasma post-treatment is performed on pre-existing plasma polymerised coatings [22]. This post-treatment, which has an effect on both the chemical composition and structure of the ppHMDSO layers,

12.1 Diagram showing the experimental atmospheric DBD plasma deposition reactor

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improves anti-corrosion passive properties. Consequently, the best deposition procedure consists of a coating step (sample exposed to plasma + precursor injection) and a post-treatment step (plasma only). The post-treatment was also 68.4 s duration. 12.2.3

Characterisation techniques

Scanning electron microscopy SEM analyses were carried out on a QUANTA 200 FEG from FEI equipped with an X GENESIS XM 4i EDS spectrometer supplied by EDAX. This microscope is a Variable Pressure SEM (VP-SEM) that enables the observation of non-conducting samples directly without additional coating [23,24]. Analyses were performed with specimen chamber pressures ranging from high vacuum (10–3 Pa) to 120 Pa depending on the charging effect. AlCeO3 powder specimens were prepared by sputtering powder onto a double-sided adhesive carbon film bonded onto an aluminium stub. Morphological investigations were performed in secondary electron mode. In order to provide better chemical contrast, the coating morphology and filler dispersion were investigated on coatings deposited onto a flat silicon substrate. Observations were performed using the backscattered electron imaging mode (BSE) and by X-ray elemental mapping (EDS). BSE analyses were carried out with a 30 kV electron probe. X-ray mapping microanalyses were carried out with a 5 kV electron probe for a live time of 900 ms per dot. Samples used to obtain general surface views were mounted directly on double-sided adhesive carbon film bonded onto an aluminium stub. To allow the observation of cross-sections, samples were broken in liquid nitrogen and held vertically in a gripping stub. Nanoscratch experiments Scratching was performed on galvanised steel samples coated with anticorrosion layers such as ppHMDSO or chromated layers. This simulates severe damage caused to these layers. Scratching experiments were carried out by a nanoscratch NST from CSM Instrument Company. A sphero-conical tip of diamond-like carbon-coated steel penetrated into the sample. All experiments, regardless of their nature (constant or increasing normal load with a constant displacement velocity of the indenter), included three successive steps performed at the same location: -

First, the profile of the original surface was determined by a scratch test at 1 mN constant normal load (prescan) The scratch test was then performed Then the residual profile of the scratch was determined by a 1 mN constant normal load scratch (postscan).

The penetration depth of the tip during the test was determined by comparing the original profile (prescan) with the vertical position of the indenter during the scratch test. Scratches of 3 μm depth were produced so that anticorrosion coatings of thicknesses much lower than 3 μm would fail. The radius of the spherical tip was 2 μm and the included angle of the conical section was 90°. A scratch with a constant normal load (50 mN) was made on samples at a scratching speed of 1 mm/min. The scratched

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length was 0.5 mm and the appearance of the residual scratch track was observed with an optical microscope after the scratch test. Interferometry experiments Three-dimensional profilometry pictures of hybrid coatings were obtained with a Wyko NT3300 white light interferometer (Veeco Company). Further topography data treatment was performed using Mountains Map software (Digital Surf Company). The coating thickness was determined by measuring the step height between the substrate surface and the deposited layer. Here the substrate under test was a glass slide, which had a very flat surface topography. Topographical grain analyses, which characterise the peak structure of ppHMDSO layers containing AlCeO3 particles, were performed using the following parameters: 0.1 μm lateral resolution (theoretical resolution), vertical shifting interferometry (VSI) mode, VSI filter, measurement size 150 μm by 150 μm, 1 measurement per sample. Data treatment consisted of filling-in non-measured points and applying a 12-order polynomial to remove the form. This high-order correction was chosen because the following steps involve working on the upper part of the AbbottFirestone curve1 to isolate the peaks clearly. It is thus of prime importance that the surrounding surface is as flat as possible. After form correction, a slight truncation of the surface was performed to remove aberrant data points (in both peak and valley areas). Peaks were then isolated using the Sr1 parameter (the 3D equivalent of Mr1 – ISO 13565-2:1996) of the Abbott-Firestone curve. This parameter, called the ‘upper material ratio’, gives the bearing ratio linked to the individual peaks. A second truncation of surface heights, from the highest peak down to the height related to Sr1, was then undertaken (Fig. 12.2). Once highlighted, peaks were counted, and their mean height, mean area and height to area ratio (i.e. peak transverse area at 15% bearing ratio) were computed. Electrochemical measurements Electrochemical experiments on galvanised steel samples were performed using a PARSTAT 2273 potentiostat/galvanostat (Princeton Applied Research) and a GAMRY 600 with a PCI4 Controller Board. Tests were carried out in 35 g/L NaCl non-deaerated aqueous solution, using a three-electrode electrochemical cell, consisting of the working electrode (1 cm2 exposed area), a saturated calomel electrode (SCE) as the reference electrode and platinum grid as the counter electrode. The measurements were performed at room temperature, during immersion of the sample in solution. The polarisation curve was measured by scanning the potential from +20 mV/SCE down to –600 mV/SCE versus the rest potential, which was the stabilised open circuit potential. This curve allowed measurement of the corrosion current, which is directly

1

The Abbott-Firestone or bearing ratio curve is the integral of the amplitude distribution function (ADF). This integral is performed from the highest peak downward. Thus, each point on this plot has the physical significance of the fraction of the projected surface above a given height, or the surface bearing ratio Ar (expressed in percent) at this height.

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12.2 A ppHMDSO coating containing 4% nanoparticles – peak heights range from white down to black. Here the Sr1 parameter is 20.1%, or 1.34 μm under the highest peak. Peak discrimination is correct

proportional to the corrosion rate, in the case of general corrosion. The EIS measurements were performed at the open circuit potential. The measurement frequency ranged from 100 kHz down to 10 mHz, with a 10 mV amplitude sinusoidal voltage. 12.3 12.3.1

Results and discussion Benchmark without AlCeO3

Anti-corrosion properties were measured using electrochemical methods: cathodic polarisation Tafel plots, which are not presented here, and Bode plots for plasma polymerised hexamethyldisiloxane ( ppHMDSO) samples (Fig. 12.3). The Tafel curves chosen correspond to the best result obtained for a series of three measurements on each sample. The open circuit potential after 3 h of immersion was taken as the reference potential (Ecorr) for each curve; hence it was possible to compare the cathodic current between different curves and therefore the anti-corrosion properties of different coatings. The corrosion current, in agreement with Tafel theory, was calculated by extrapolating the oxygen reduction plateau at the corrosion potential [25,26]. In comparison with the untreated substrate, a significant decrease in the corrosion current was observed for ppHMDSO-coated samples. Indeed, the average value corresponding to these siloxane films was about 1.2×10–7 A/cm2 whereas the aluzinc sample gave a cathodic current of 3.5×10–7 A/cm2. As a reference, the curve for the chromated coating gave a value of 5.0×10–8 A/cm2. The plasma polymerised coatings improved the corrosion resistance by a factor of three compared to the untreated sample, and the current value was close to the value for chromated coatings. The corrosion resistance of these coatings is due to the development of a dense –Si–O–Si– network, which

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12.3 EIS Bode plots of galvanised steel substrates, chromated and ppHMDSO layers recorded after 50 h of immersion

hinders the penetration of aggressive species towards the metallic substrate. The corrosion protection effectiveness of siloxane coatings is strongly dependent on their barrier properties. Improvement of the barrier properties of siloxane coatings is a prominent objective towards the development of corrosion protection layers. The impedance spectra can be investigated to allow modelling of the physicochemical processes that occur in the coating–substrate system during corrosion tests. Figure 12.3 shows the evolution of the impedance spectrum of ppHMDSO coating with, for reference purposes, spectra from untreated and chromated samples. Experiments are commenced after 50 h of immersion in NaCl solution. Concerning the Z modulus curves, at high frequencies, chromated and ppHMDSO coatings display the same characteristics. At low frequencies, slightly better corrosion resistance is observed for the chromated film (300 kΩ) than the polymer one (100 kΩ). On the other hand, the resistance of the untreated substrate was lower by one order of magnitude. Therefore the results obtained from Tafel plots and impedance spectra are consistent. 12.3.2

Addition of AlCeO3 nanoparticles

Morphological investigation Elemental cartography and topography Figure 12.4a shows a general view at low magnification of the surface of a ppHMDSO coating containing 3% AlCeO3 nanoparticles with a thickness of 300 nm. The coating surface is neither smooth nor flat. Numerous irregularities can be seen,

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especially protuberances that appear white in the micrograph. These protuberances were present on the whole surface and were uniformly distributed in the matrix. Their diameter ranged from a few hundred nanometres to a few micrometres. Aluminium elemental mapping (Fig. 12.4b) showed clearly that these protuberances correspond to Ce-based particles, which were the only constituents containing aluminium. Since the dimensions of these structures were of micrometre magnitude, the AlCeO3 particles must have formed agglomerates that were distributed homogeneously in the matrix. The elemental mapping for silicon (Fig. 12.4c) confirmed this observation. The silicon signal was almost undetectable at positions corresponding to nanoparticle agglomerates. It came from the substrate and the ppHMDSO coating. Elemental mapping for oxygen (Fig. 12.4d) was perfectly consistent with these observations. The most intense signals correspond to the particles, the signals of medium intensity to the extra thickness around the particles, and the weakest signals to the matrix background. It can be noted that the darkest areas of oxygen mapping correspond to the masking effect of the largest agglomerates. To confirm these observations, a thorough study was undertaken on both the sample surface and cross-sections (Fig. 12.5). Particles are gathered into agglomerates whose diameter ranges from 100 nm to several micrometres. It was noticed that similar agglomerates were present in the initial powder. Figure 12.5b, obtained on a cross-section with a tilt of 5°, confirmed that particles were embedded in the polymer matrix. Agglomerates whose diameter was greater than 200 nm (the thickness of the coating) were thicker than the surrounding film. Indeed, an excess thickness of 150 nm relative to the agglomerate is visible in this image. Figure 12.5b1 represents agglomerates in the BSE imaging mode, where nanoparticles appear white and are embedded in the matrix.

12.4 SEM micrograph (a) and mapping scan of Al (b), Si (c) and O (d) of a ppHMDSO coating containing 3% AlCeO3 nanoparticles

12.5 Backscattering SEM micrograph of a ppHMDSO coating containing 3% AlCeO3 nanoparticles: (a) surface; (b) cross-section; (b1) BSE; (b2) SE

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Influence of nanoparticle concentration The influence of particle concentration on the coating structure and morphology was studied by correlation of SEM results and interferometer analyses. Samples with nanoparticle concentrations of 1% to 4% were analysed. Figure 12.6 shows the microscopy results. The SEM micrographs correspond to an area of 250 × 250 μm. Particle agglomerates appear in white because they are composed of heavier elements than the plasma polymer. These micrographs enable evaluation of the structure and distribution of the smallest agglomerates as a function of the concentration of particles. The number of observable agglomerates increases with the concentration of particles. These hybrid coatings were also studied by interferometry: mean values are reported in Table 12.1. Figure 12.6(f to j) shows the full three-dimensional measurement before peak isolation. Peaks become more visible as the concentration of nanoparticles is increased, i.e. they become more numerous and higher (especially for a concentration of 4%). Since the reference state (no particles) is free of peaks (Fig. 12.6f), these peaks can be linked directly to the presence of cerium particles. The numerical values from Table 12.1 confirm this visual trend. When the three other samples were studied further, calculations showed that the peaks became more numerous with increasing nanoparticle concentration, without significant change in both mean height or area value, but at higher concentration levels, they became significantly higher and thinner (see height over surface ratio). These observations are in good agreement with SEM measurements (same order of magnitude of peaks in both techniques, see Fig. 12.5). Further topographical analyses of peaks will be carried out. Electrochemical measurements Influence of nanoparticles Plasma polymerisation of coatings incorporating AlCeO3 nanoparticle concentrations in the range 1 to 4% was performed. Tafel plots corresponding to these films are presented in Fig. 12.7a in comparison with chromated coatings. The curves are characterised by a current plateau relative to the limiting current of dissolved oxygen; depending on the conditions, this cathodic process can be explained by two different mechanisms, as described by equations 12.1 and 12.2, respectively: O2 + 2H2O + 2e– Æ H2O2 + 2OH–

[12.1]

O2 + 2H2O + 4e Æ 4OH

[12.2]

–

–

This phenomenon is usually observed for experiments carried out on bare metals immersed in corrosive solutions, showing that the corrosion mechanism is under diffusion control [26,27]. In the –0.1 to 0 V range, corresponding to dioxygen reduction, corrosion currents relative to filled samples decrease slowly and range between 9.0×10–8 and 1.2×10–7 A/cm2. The presence of nanoparticles improves corrosion resistance when the concentration of particles is greater than 2%. The main anodic reaction during the corrosion of untreated samples is oxidation of aluminium and zinc. At cathodic sites, due to the local pH increase (equations 12.1 and 12.2), the deposition of cerium(III)

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12.6 SEM micrographs and 3D interferometry pictures of ppHMDSO containing AlCeO3 at concentrations of 0% (a, f), 1% (b, g), 2% (c, h), 3% (d, i) and 4% (e, j)

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Table 12.1 Influence of nanoparticle concentration on article characteristics

Number of peaks Mean height Mean area Height/area ratio

[nm] [μm2] [nm/μm2]

1% of particles

2% of particles

3% of particles

4% of particles

232 78 6.4 12.0

294 40 4.2 9.6

293 74 4.8 15.4

319 198 12.0 16.5

hydroxide may occur, blocking the cathodic reaction and hindering the whole corrosion process. EIS spectra measured on these coatings are presented in Fig. 12.7b. These impedance values confirm that the barrier properties, visible at low frequencies, are improved by AlCeO3 incorporation in comparison with the reference sample. The best corrosion resistance is achieved when the concentration of particles is 3%: beyond this value, it seems that further incorporation does not improve corrosion resistance. The incorporation of this kind of particles enhanced the ppHMDSO barrier properties compared to those of pure ppHMDSO films. The corrosion resistance results are consistent between the Tafel plots and Bode plots. In both cases, the best corrosion resistance is obtained with 3% incorporation. Immersion time During immersion, defects appear in the protective coating due to water uptake and they become preferential sites for corrosion. In order to see the evolution of these defects with the different kinds of coating, the films were immersed in the aggressive solution for a long time. Figure 12.8 shows the evolution of the Z modulus for pure and AlCeO3-filled ppHMDSO films. The coatings were immersed in NaCl solutions for 5 days. After this period, the good barrier properties of ppHMDSO films decrease. Similarly, the global impedance values decrease by 80 kΩ to 40 kΩ. On the other hand, curves corresponding to ppHMDSO containing AlCeO3 nanoparticles are more or less superimposed, whatever the immersion time. Passive properties (barrier effect) and global values of Z are similar. The addition of corrosion inhibitor to the siloxane matrix can confer additional active corrosion protection. Cerium salts can provide self-healing ability to supplement the good barrier properties of the films. Corrosion inhibition at localised defects In order to investigate the self-healing ability of different siloxane coatings, artificial defects were made in the surface of different films. When the barrier coating is damaged, corrosive species can reach the metal substrate, leading to corrosion activity. Thus, the presence of an active inhibiting component in the protective coating helps to lower corrosion activity. The main goal of cerium additions to the siloxane coating is to impart active corrosion protection properties to the ppHMDSO barrier layer. Active corrosion protection means that the sample can hinder corrosion activity due to damage of the barrier layer at the defects. Therefore, artificial defects were created in the siloxane coatings to determine how hybrid films behave when the barrier is damaged.

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12.7 Tafel plots (a) and impedance plots (b) of ppHMDSO layers containing AlCeO3 nanoparticles

Figure 12.9 presents impedance spectra obtained on chromated (a), ppHMDSO (b) and ppHMDSO containing 3% AlCeO3 (c) samples, in which an artificial defect was made by the introduction of a nanoscratch after 1 day of immersion. The samples were first immersed in NaCl solution for 24 h. After this period, a defect was created on the surface and new EIS spectra were measured. According to the literature, two kinds of behaviour can be expected. In the first [28], a decrease in impedance values after the creation of the defect is observed, but after more than 24 h of immersion,

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12.8 Evolution with immersion time of impedance plots of ppHMDSO layers with and without AlCeO3 nanoparticles; 1d, 1-day immersion in NaCl solution; 5d, 5-day immersion in NaCl solution

this impedance remains approximately constant: the presence of active nanoparticles blocks further evolution of corrosion. In the second behaviour, a decrease in the Z value occurs after the scratch, followed by a slight increase in impedance at low frequencies, revealing that corrosion has slowed down [29]. Chromated coatings are well known for their self-healing properties, and the impedance spectra (Fig. 12.9a) confirm this behaviour: whatever the time of immersion or presence of a scratch, the curves can be superimposed. A slow decrease in the global resistance is observed after 3 h, then stabilisation occurs: chromate ions play their inhibiting effect and stop corrosion activity. With the pure ppHMDSO layer (Fig. 12.9b), the global corrosion resistance decreases slightly 3 h after the creation of the defect, and then more strongly after 24 h. Corrosion activity increases when the metal is in contact with an aggressive solution, which is consistent with the evolution of the curve. The siloxane-based polymer does not provide any active protection. On the other hand, when AlCeO3 particles are incorporated, a decrease in the global corrosion resistance is observed immediately after the scratch (Fig. 12.9c). However, after longer immersion, the resistance increases strongly and tends to the initial value. The anticorrosive effect of the cerium ions entrapped into the siloxane coating is due to an inhibiting effect and a self-healing mechanism, probably with cerium hydroxide precipitation. Indeed, the inhibitive action of Ce3+ is based on the formation of highly insoluble hydroxides, which can be formed in the place where the cathodic reaction occurs according to the following equation: Ce3+ + 3OH– Æ Ce(OH)3

[12.3]

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12.9 Impedance spectra of chromated (a), ppHMDSO (b) and ppHMDSO layer containing 3% AlCeO3 (c) – one scratch

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Indeed, when nanoparticles are embedded in the siloxane coating, the following mechanism can be assumed: the silica network progressively breaks down, which releases the nanoparticles that precipitate on the electrode surface. Hence, they form complexes with charged species, which reinforces their protective role. The formation of these stable corrosion products decreases the active corrosion area, hence slowing down the corrosion activity. 12.3.3

Optimisation of active properties

In order to improve our understanding of the active properties of coatings containing Ce-based nanoparticles, two kinds of experiments were undertaken. On the one hand, more scratches were made on the best coating, i.e. ppHMDSO containing AlCeO3, and on the other hand, an attempt was made to improve the dispersion of the nanoparticles. The EIS spectra resulting from these tests were studied. Increase in localised defects Figure 12.10 shows results obtained for samples scratched five times: this corresponds to five times more scratches than in the previous experiments (Fig. 12.9). A slight decrease in the global corrosion resistance was observed 1 day after immersion and 3 h after creation of the defects. However, after immersion for 24 h following the introduction of the scratches, this decrease stopped and the curves became superimposed. The evolution of the global impedance (in percent of the arbitrary best Z value obtained after 3 h of immersion) was calculated and results are presented in Table 12.2.

12.10 Impedance spectra of ppHMDSO layer containing 3% AlCeO3 – 5 scratches

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Table 12.2 Variation of impedance with time as a percentage of the maximum value (3h before scratching) Chromated

3 h before scratch 24 h before scratch 3 h after scratch 24 h after scratch

1 scratch

ppHMDSO +AlCeO3 3% 1 scratch

ppHMDSO ppHMDSO +AlCeO3 3% +AlCeO3 3% (EtOH) 5 scratches 1 scratch

Z max = 100 67% 59% 59%

Z max = 100 92% 36% 54%

Z max = 100 76% 64% 59%

Z max = 100 98% 84% 79%

Comparison shows that almost 60% of Zmax remains after scratching the chromated sample, and this value is constant after 24 h of immersion. For ppHMDSO containing AlCeO3, this percentage decreases strongly to 36% immediately after the introduction of one scratch then moves up to 54% thanks to its self-healing properties. On the other hand, the evolution of this kind of coating in relation to five scratches is different. The global impedance decreases to 64% after scratching and stabilises at around 60%. In spite of the greater number of defects on the surface, the corrosion does not increase as much as might be expected. It seems that a sufficiently large release of cerium ions occurred to slow down the corrosion reactions. In conclusion, more scratches result in a more pronounced self-healing reaction: there is not an increase in Z value as observed with one scratch but a tailing off of Z decrease. Improvement of nanoparticle dispersion Self-healing properties depend on cerium ions and their ability to be released into the polymer matrix. In order to optimise the dispersion of nanoparticles, a small amount of ethanol was added to the initial solution of HMDSO and AlCeO3 (3%) powder. Figure 12.11 presents SEM results for this surface, in comparison with those obtained without solvent. Some thick aggregates, which are present in Fig. 12.11a, disappear when the particles are dispersed in ethanol. Protuberances are less evident in this case and it seems that nanoparticles do not agglomerate. Concerning anti-corrosion properties, Tafel curves allow the calculation that the corrosion current for this kind of coating is slightly below that for classical Ce-based ppHMDSO films. Images of HMDSO treated samples after 25 days in a salt spray chamber are shown in Fig. 12.11(c,d) which compare the resistance of HMDSO plasma treatments with and without ethanol. An increase in corrosion protection is afforded by the presence of solvent, since pitting corrosion is reduced. This last result corroborates those obtained by electrochemistry. Active properties have also been studied using the same protocol as previously. The results are presented in Table 12.2. After 24 h of immersion in corrosive solution, 98% of the maximum impedance is retained, which is better than for the chromated sample or the (ppHMDSO+AlCeO3) coating. The same is true of values obtained after scratching, whatever the time of immersion. Indeed, the global impedance decreases extremely slowly and is close to 80% of maximum of Z after 1 day of immersion. AlCeO3 nanoparticles seem more dispersed into the polymer matrix and slow down corrosion activity thanks to good self-healing properties.

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12.11 SEM micrographs and salt spray test images of ppHMDSO containing AlCeO3 (3%) without (a, c) and with ethanol (b, d)

12.4

Conclusion

Plasma polymerised HMDSO coatings containing AlCeO3 nanoparticles and deposited on galvanised steel by atmospheric pressure dielectric barrier discharge have been studied. Morphological investigations have confirmed that nanoparticles are incorporated into the polymer matrix with an effective distribution, whatever the concentration. In addition, an excess thickness of these particles within the matrix (protuberances) is sometimes visible. This phenomenon has been confirmed by interferometry measurements, which show a surface morphology with a number of peaks that is proportional to the concentration of nanoparticles. Electrochemical experiments have confirmed the good barrier properties of hybrid siloxane coatings and also the active corrosion protection effect of cerium ions. Indeed, polymer coatings with cerium incorporation performed at least as well as the undoped siloxane layers: the best barrier properties were obtained with a 3% AlCeO3 concentration. Concerning the self-healing behaviour, ppHMDSO coatings incorporating cerium nanoparticles provide improved long-term corrosion protection and active properties.

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Acknowledgements The authors would like to thank Mr Fiorucci and Mr Kloos (Arcelor Mittal Dudelange) for constructive discussions and for providing galvanised steel samples. This work (TRASU project) is sponsored by the Luxembourg research funding association FNR (Fond National de la Recherche). References 1. http://www.epa.gov/safewater/contaminants/dw_contamfs/chromium.html. Date last accessed 2. E.W. Brooman, Met. Finish., 100(5) (2002), 42–53. 3. E.W. Brooman, Met. Finish., 100(1) (2002), 48–53. 4. E.W. Brooman, Met. Finish., 100(6) (2002), 104–110. 5. M. F. Montemor, A. M. P. Simoes, M. G. S. Ferreira, B. Williams and H. Edwards, Prog. Org. Coat., 38 (2000), 17–26. 6. M. F. Montemor, A. M. Cabral, M. L Zheludkevich and M. G. S. Ferreira, Surf. Coat. Technol., 200 (2006), 2875–2885. 7. J. Bardon, J. Bour, H. Aubriet, D. Ruch, B. Verheyde, R. Dams, S. Paulussen, R. Rego and D. Vangeneugden, Plasma Proc. Polym., 4 (2007), 445–449. 8. T. P. Chou, C. Chandrasekaran, S. J. Limmer, S. Seraji, Y. Wu, M. J. Forbess, C. Nguyen and G. Z. Cao, J. Non-Cryst. Solids, 290 (2001), 153–162. 9. V. Palanivel, Y. Huang and W. J. Van Ooij, Prog. Org. Coat., 53 (2005), 153–168. 10. M. Sheffer, A. Groysman, D. Starosvetsky, N. Savchenko and D. Mandler, Corros. Sci., 46 (2004), 2975–2985. 11. A. Pepe, M. Aparicio, S. Ceré and A. Duran, J. Non-Cryst. Solids, 348 (2004), 162–171. 12. A. M. Cabral, W. Trabelsi, R. Serra, M. F. Montemor, M. L. Zheludkevich and M. G. S. Ferreira, Corros. Sci., 48 (2006), 3740–3758. 13. M. F. Montemor, W. Trabesli, M. Zheludkevich and M. G. S. Ferreira, Prog. Org. Coat., 57 (2006), 67–77. 14. V. Palanivel, D. Zhu and W. J. Van Ooij, Prog. Org. Coat., 47 (2003), 384–392. 15. C. Nastase, F. Nastase, A. Dumitru, M. Ionescu and I. Stamatin, Composites, 36 (2005), 481–485. 16. C. Nastase, F. Nastase, A. Vaseashta and I. Stamatin, Prog. Solid State Chem., 34 (2006), 181–189. 17. F. Nastase, I. Stamatin, C. Nastase, D. Mihaiescu and A. Moldovan, Prog. Solid State Chem., 34 (2006), 191–199. 18. M. Dabala, E. Ramous and M. Magrini, Mater. Corros., 55 (2004), 381–390. 19. F. Mansfeld, J. Electrochem., 36 (2000), 1063–1069. 20. C. B. Breslin, C. Chen and F. Mansfeld, Corros. Sci., 39 (1997), 1061–1073. 21. A. S. Hamdy, Surf. Coat. Technol., 200 (2006), 3786–3792. 22. J. Bour, J. Bardon, H. Aubriet, D. Del Frari, B. Verheyde, R. Dams, D. Vangeneugden and D. Ruch, Plasma Proc. Polym., 5 (2008), DOI: 10.1002. 23. G. D. Danilatos, Microsc. Res. Tech., 25 (1993), 529–534. 24. V. N. E. Robinson, J. Phys., 8 (1975), 638–640. 25. C. Fiaud, S. Bensara, I. Demesy des Aulnois and M. Tzinmann, Br. Corros., 22 (1987), 102–108. 26. S. Manov, A. M. Lamazouère and L. Aries, Corros. Sci., 42 (2000), 1235–1248. 27. C. Deslouis, M. Duprat and C. Tulet-Tournillon, J. Electroanal. Chem., 181 (1984), 119–136. 28. W. Trabelsi, P. Cecilio, M. G. S. Ferreira and M. F. Montemor, Prog. Org. Coat., 54 (2005), 276–284. 29. M. L. Zheludkevich, K. A. Yasakau, A. C. Bastos, O. V. Karavai and M. G. S. Ferreira, Electrochem. Commun., 9 (2007), 2622–2628.

13 Polypyrrole/aluminium flake hybrid pigments for corrosion inhibition of aluminium 2024

Chris Vetter, Xiaoning Qi, Subramanyam V. Kasisomayajula and Victoria Johnston Gelling Department of Coatings and Polymeric Materials, North Dakota State University, 1735 NDSU Research Park Drive, Fargo, ND 58105, USA [email protected]

Alice C. Harper Department of Chemistry, Berry College, 2277 Martha Berry Hwy. NW, Mount Berry, GA 30149, USA

13.1

Introduction

With the paper ‘Synthesis of electrically conducting organic polymers: Halogen derivatives of polyacetylene (CH)x’ in 1977, the first research into electroactive conducting polymers such as polypyrrole was started [1]. With a high conductivity of 105 Siemens per metre, the doped polyacetylene had a conductivity which was 109 times more conductive than that of the undoped polyacetylene and was also on a par with the conductivity observed for many metals. It was determined that the crucial factor determining the conductivity of the polymer was the conjugated double bonds at the backbone of the conducting polymer. Professor Alan J. Heeger, Professor Alan G. MacDiarmid, and Professor Hideki Shirakawa were awarded the 2000 Nobel Prize in Chemistry in part for the publication of this paper and the resulting research [2]. In 1985, the first work utilising an organic conducting polymer as a corrosion inhibitor was published by DeBerry. The study probed coating stainless steel electrodes with polyaniline thin films and subsequent exposure to an acidic solution. The results of his work showed that the coated electrodes were passivated in acidic solutions while the same electrodes that were left untreated were more rapidly corroded [3]. During the late 1990s, the amount of literature being produced on the synthesis and applications of organic conducting polymers permitted the publication of several extensive review articles. Reviews have been authored by McAndrew et al., Stoffer et al., Yagova et al., Lu et al., Tallman et al., and Spinks et al. that address the use of electroactive conducting polymers (ECPs) as corrosion inhibitors [4–9]. There are several proposed mechanisms for the corrosion inhibiting behaviour observed for polypyrrole. One mechanism proposed in the literature is the release of a corrosion inhibiting counter ion from the polypyrrole as it is reduced from its oxidised state to its neutral state by the electrons produced by the oxidation of the metal substrate. This mechanism can be thought of as a ‘smart’ or responsive coating 238

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that is able to detect the onset of corrosion and release inhibitors in a controlled manner. Indeed, Paliwoda-Porebska et al. utilised a Scanning Kelvin Probe to study the release of anions from a conducting polymer [10]. Another proposed mechanism by which polypyrrole may inhibit corrosion relies upon its inherently positive open circuit potential when polypyrrole (PPy) is electrically coupled with a metal substrate. Tallman et al. suggested that PPy joined to aluminium results in a coupled potential that is more positive and therefore more noble than that of the uncoupled metal [11]. A significant amount of research has been performed to increase the practicality of using ECPs on a large scale by increasing the solubility and processability. The basis for this research is the fact that polypyrrole is not soluble in most common solvents. To synthesise a polypyrrole that is soluble in organic solvents, alkyl substituents can be added to the polymeric backbone at the β position on the pyrrole ring. The synthesis process yielding alkyl-substituted pyrrole has produced a polymer that is soluble in several organic solvents such as carbon tetrachloride, dichloromethane, and acetonitrile [12–14]. Due to the environmental concerns of using solvent-borne synthesis routes, the need for an aqueous synthesis process is desirable. Many groups have been investigating the aqueous electropolymerisation of pyrrole onto surfaces. This is made possible by the pyrrole monomer’s modest solubility in water (~1 M) [8,15–17]. This is a limited technique, however, as it can only be used for conductive substrates [15]. To overcome the necessity of having a conductive substrate, it is necessary to utilise the chemical oxidation of pyrrole. In this method, strong chemical oxidants, such as ammonium persulphate, are used to initiate the polymerisation reaction. The resulting product can be used as is or used as a pigment when combined with other coating technologies. This strategy may incorporate the superior mechanical properties of the coating binder with the corrosion inhibiting properties of the polypyrrole. Several research groups have used PPy as an additive into corrosion inhibitive coatings. Armelin and colleagues performed a study on the use of polyaniline and polypyrrole as additives incorporated into an epoxy paint [18,19]. In the studies, the conducting polymers were purchased from commercial sources and incorporated at relatively low concentrations into an epoxy paint. The corrosion assessment consisted primarily of accelerated cabinet exposure with subsequent visual inspection. The results indicated that polyaniline and polypyrrole were able to provide shortterm corrosion inhibition as detected by a decrease in the spread of corrosion from a defect site. A similar approach was undertaken by Amarnath et al., Aradilla et al., and Sathiyanarayanan et al. All incorporated conducting polymers into traditional barrier-type coatings at varying concentrations and studied the resulting corrosion performance of the coating [20–22]. One solution to mitigate the undesirable properties of polypyrrole is to directly deposit polypyrrole onto aluminium flakes and incorporate the resulting hybrid pigment at various pigment volume concentrations (PVCs) into an epoxy primer. In comparison to the multiple studies on the use of conducting polymers used as corrosion inhibiting additives, the use of conducting polymer composites for use as anti-corrosion additives is not as thoroughly investigated. Upon review of the literature, it was found that Sathiyanarayanan and colleagues have completed various studies on polyaniline composites including polyaniline/Fe2O3, polyaniline/glass, and polyaniline/TiO2 [23–25]. In these studies, the polyaniline composites were incor-

Polypyrrole/aluminium flake hybrid pigments for corrosion inhibition

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porated into polymeric resins; however, it should be noted that other pigments such as TiO2 or mica were included in many of the formulations. The samples were exposed to a corrosive electrolyte and electrochemical impedance spectroscopy (EIS) was performed. Interestingly, the polyaniline/glass (total PVC ~16%) and the polyaniline/TiO2 (total PVC ~35%) composite samples both displayed relatively high impedances, indicative of a barrier-type coating below the critical pigment volume concentration or percolation threshold for a conductive pigment. The Fe2O3 composite (total PVC ~35%), in contrast, displayed relatively low impedances. This would seem to indicate that there was a difference in conductivities between the TiO2 and Fe2O3 composites as both were formulated at similar PVCs and yet displayed extremely different behaviours when examined by EIS. The interest in polypyrrole/aluminium flake composite is twofold. The aluminium flake can act as a filler to decrease the amount of polypyrrole needed to form a conductive coating and may impart some barrier properties to the coating by creating a difficult pathway for ions and water to diffuse through the coating. In a coating such as this, the surface morphology and amount of polypyrrole that is deposited onto the aluminium flake would be a critical factor affecting the properties of the coating. If the flakes are not sufficiently covered with PPy, it would be difficult to form electrical connections with the substrate and throughout the coating thereby decreasing the ability of the coating to inhibit corrosion. Upon review of the literature, it was found that catechol had been used during the electrodeposition of polypyrrole on aluminium substrates. The catechol lowered the potential required for deposition and resulted in a more uniform coating [26]. In this study, the effects of catechol, a phenol, on the resulting composite pigment formed via chemical oxidation were studied. Additionally, other phenols were used to determine the optimum experimental conditions to produce a uniform deposition product on the aluminium flake. 13.2

Experimental

Most chemicals were used as received such as catechol (TCI America); ammonium persulphate, phloroglucide, and phloroglucinol (Sigma–Aldrich); ethanol and tetrachloroethylene (J.T. Baker); epon 828 and epikure 3175 (Hexion); methyl isobutyl ketone (Aqua Solutions); resorcinol (MP Biomedicals); and methyl ethyl ketone (Alfa Aesar). The aluminium flake used was an automotive pigment, Stapa Aloxal PM 2010 (Eckhart America). The pyrrole monomer (TCI Amercia) was distilled before use. The procedure for the synthesis of the polypyrrole/aluminium flake composite pigment proceeded as follows (as detailed in Table 13.1 and Fig. 13.1): 3 g of aluminium flake was dispersed into 100 ml of 18.2 MŲ Millipore water. To this, ammonium persulphate (APS) was added, followed by one or more phenolic compounds and then pyrrole was added. This mixture was allowed to stir for 24 h at which point the product was filtered and washed with 150 mL of 18.2 MŲ Millipore water. The PPy/Al flake composite was then placed in an oven and allowed to dry for 24 h at 60°C. Once the PPy/Al flake composite had dried, it was ground using a mortar and pestle and passed through a 150 μm sieve, which was obtained from VWR. The ground composite pigment was then incorporated into an epoxy primer at various pigment volume concentrations (PVC). The primer was prepared by mixing epon 828 and epikure 3175 in a 1:1.2 ratio and diluting it with methyl isobutyl ketone. The composite pigment was then added and more methyl isobutyl ketone was added

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Self-healing properties of new surface treatments

Table 13.1 Synthesis reactions for PPy/Al flake hybrid pigments. Shaded area denotes lack of chemical in the synthesis Reaction Catechol Phloroglucide Phloroglucinol Resorcinol Pyrrole Phenol + Pyrrole # (M) (M) (M) (M) (M) concentration (M) 1 1x 2 3 4 4x 5 6

0.1 0.2 0.1

0.1 0.1 0.1 0.2

0.1

0.1 0.1

0.1 0.2 0.2 0.1 0.1 0.2 0.2 0.1

0.2 0.4 0.4 0.2 0.2 0.4 0.4 0.2

13.1 Reactants used in the PPy/Al flake synthesis

with stirring until a proper application viscosity was reached. The primer was then shaken vigorously by hand for approximately 10 min. The panels used as substrates were supplied by Q-Panel and were composed of aluminium alloy 2024 T3. Before coating the panels with the composite epoxy primer, they were sanded using 160-grit sandpaper and a Master Mechanic® palm sander. The panels were then degreased using ethanol. While this is not a standard industrial protocol for aluminium substrates, it was performed to degrease and clean the surface while reducing any corrosion inhibition that would be provided by a standard industrial cleaning and pretreatment step. Samples that were to undergo ASTM B117 salt spray exposure (B117) were applied using an ATD Air-6900 high volume low-pressure spray gun with a 1.8 mm orifice. All other coatings were applied using a drawdown bar supplied by Gardco at 8 mils wet thickness. The panels were allowed to flash off for 40 min before being placed in an oven at 80°C to cure for 4 h. The final cured thickness of the primers was from 127 μm to 178 μm. Half of the samples that were to undergo B117 exposure were scribed using a Gravograph IM4 engraver. A PHI model 555 XPS instrument was used to perform XPS measurements using Mg K-alpha X-ray excitation with a cylindrical mirror electron energy analyser. A pass energy of 100 eV was used to collect survey photoelectron spectra. Highresolution photoelectron spectra were collected using a 25 eV pass energy. To neutralise the samples during measurements, a low energy flux (1.0 eV, 0.9 mA) was

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applied. Charge compensations were done, during curve fittings, by setting the C 1s peaks for C–C/H at 285.0 eV, and a Gaussian-Lorentzian model with a 90:10 ratio was applied. To prepare samples for SEM, the ground composite pigment was sprinkled onto carbon tape which was attached to aluminium mounts. Compressed air was blown onto the sample to remove any loose particles and the sample was then sputtercoated with gold using a Balzers SCD 030 sputter coater. The instrument used to obtain images was a JEOL JSM-6300 Scanning Electron Microscope. The magnification, accelerating voltage used and scale bars are listed on each figure. Ground product was used for the conductive AFM (C-AFM) measurements. The sample substrates were prepared by first sanding a small aluminium panel and then applying some conductive silver epoxy to the surface. The sample was dusted onto the surface of the epoxy. This was allowed to cure overnight before measurements. The samples were measured using a Digital Instruments Dimension 3100 AFM from Veeco. The samples were measured in contact mode using platinum/iridium coated tips from Veeco. All samples were first scanned with no applied voltage to ensure a good image could be obtained. After this was verified, a voltage of 100 mV was applied to the sample and the resulting current image was captured. The conductivity was calculated by using the following equation: conductivity = (measured current/applied voltage) × (sample thickness/contact area of tip)

[13.1]

2

The contact area of the tip used was calculated to be 19.5 nm . The sample thickness was determined from the cross section of the height image, and the average current was determined from the cross section of the current image. Electrochemical impedance spectroscopy (EIS) was used to monitor the samples in B117 exposure where periodically the samples were removed from the chamber and EIS measurements were taken from non-scribed areas of the panels. EIS is a method where a small voltage perturbation is applied to a sample, usually about its open circuit potential. The current flow due to the voltage perturbation is measured and a form of Ohms law is used to calculate the impedance (ohms) as the voltage and current are known: Z (v) =

V (v) I (v)

[13.2]

Z(ƒ) is the complex impedance which includes the relationship between amplitudes of the voltage and current signals as well as any phase shift that may occur. V(ƒ) and I(ƒ) are the voltage perturbation and corresponding current, respectively. EIS was performed using a Gamry PC-4 or FAS1 instrument with dilute Harrison’s Solution (DHS) (0.35% (NH4)2SO4 and 0.05% NaCl in 18.2 MŲ Millipore water) as the electrolyte, platinum-coated mesh as the counter electrode, and a saturated calomel reference electrode. A 10 mV perturbation was used with a frequency range from 100 000 to 0.01 Hz. The EIS cell consisted of a glass tube with rubber O-ring to seal the tube onto the coating surface. The area of electrolyte contact on the coating surface within this cell was 7 cm2. The same experimental set-up was used to collect the open circuit potential (OCP) of the panels. To perform the density tests, the ground sample was placed in tetrachloroethylene. The samples were given sufficient time to allow the higher density material to settle

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Table 13.2 Experiments performed for each reaction. Shaded area denotes that the experiment was not performed for the reaction Reaction #

SEM

Density

XPS

C-AFM

OCP

B117

EIS

1 1x 2 3 4 4x 5 6

out of the liquid at which time a digital photograph was taken and the results recorded. Refer to Table 13.2 for a summary of experimental procedures performed for each experimental reaction. 13.3 13.3.1

Results and discussion SEM characterisation of polypyrrole flakes from reactions 1–6

As can be seen in Fig. 13.2, the as-received aluminium flakes are reasonably uniform in size ranging from approximately 10 to 30 μm in diameter. According to the MSDS supplied with the flake by the manufacturer, it is composed of aluminium, aluminium oxide, and 1-methoxy-2-propanol; a small amount of impurities was also present. It is evident that the flake-like nature of the pigment was not significantly compromised during the synthesis process. However, in some instances, such as reactions 1 and 1x, large particles are present which are clumps of multiple flakes. It is interesting to note that this only occurs when catechol is the only phenol present during the reaction. The results from reaction 1 are presented in Fig. 13.3 where the deposited product on the flakes appears to join flakes together. In reaction 1x, the effect of increasing the reactant concentration in the flake synthesis was studied. The product of this reaction can be seen in Fig. 13.4. While it appears that the majority of the aluminium flakes are covered with polypyrrole, the clumping seen in reaction 1 has been amplified resulting in still larger particle sizes. This effect presented a problem in the processing of the pigment, as it was exceptionally hard to grind it to a size that would pass through a 150 μm sieve. Rather uniform deposition resulted from the combination of catechol and phloroglucide during reaction 2. As can be seen in Fig. 13.5, this reaction yielded flakes which were apparently uniformly coated. The uniform coating of the flake would most likely facilitate the formation of a conductive epoxy coating. Additionally, the flakes appear to be rather individual in nature when compared with the aforementioned reaction 1. This is particularly interesting as the overall concentration of the reactants is the same between reaction 1x and 2, the only difference being the presence of phloroglucide. Figures 13.6 and 13.7, reactions 3 and 4, also show unique polymer growth. In Fig. 13.6, when the phloroglucide is the lone phenol, such as in reaction 3, there is the

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13.2 SEM micrograph of the as-received aluminium flake (top) and the composition (from MSDS) (bottom)

development of lamellar folds of PPy similar in appearance to crumpled tissue paper. In Fig. 13.7, when phloroglucinol is the lone phenol, the flakes appear to resemble the as-received flakes with no clumping observed. However, on close examination, it appears that the product is rather unevenly distributed on the surface of the flake. It appears that when phloroglucinol is the lone phenol present during the reaction, it tends to direct the polymer growth to the edges of the flake with some wire growth towards the middle. Interestingly, when the concentration of reactants used in reaction 4 is increased in reaction 4x, as shown in Fig. 13.8, polymer deposition on the centre of the flakes is not facilitated. It simply increases the amount deposited on the edges. This is not necessarily desirable for a pigment application. The structures extending from the edges of the flakes would readily break off during the formulating and mixing of a paint. The increase in the free product would drastically influence the pigment volume concentration. The use of catechol combined with resorcinol, reaction 5, can be seen in Fig. 13.9. It was noticed that the joining of multiple flakes together as observed for reactions

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13.3 SEM micrographs of ground product from reaction 1, at original magnifications of x 550 (top) and x 5,000 (bottom)

1 and 3 was diminished by the addition of resorcinol. This is similar to the result observed for the addition of phloroglucide. While there is some product deposited from this reaction, it is not as uniform and occurs on fewer flakes than is observed in reaction 2. In Fig. 13.10, there is the presence of wire-like structures on the micro scale. These structures are also observed in Figs. 13.7 and 13.8 although they are not as prominent as in Fig. 13.10. Because of the sputtered gold coating placed on the samples for SEM measurements, it is impossible to discern the actual thickness of these wires. The synthesis of these PPy wire structures has been reported by other groups [27–29]. There are presently three ways that these structures have been synthesised. One is by using surfactants to control the polymerisation process. Another is by using templates to guide the polymerisation as it takes place and a third is by using strictly controlled electrochemical techniques to grow PPy nanowires. While there is not a traditional surfactant present during the synthesis, the asreceived flakes are prepared with a significant amount of 1-methoxy-2-propanol

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13.4 SEM micrographs of ground product from reaction 1x, at original magnifications of x 550 (top) and x 5,500 (bottom)

(35%) present to encourage easy dispersion in aqueous paints. It may be possible that the 1-methoxy-2-propanol is affecting the deposition of the polypyrrole. However, if that were the case, one would assume that the wire-like structures would appear on the products for every reaction, which is not the case. Certainly, differences in the resulting morphology of the deposited product appear linked to the phenolic compounds used during the synthesis. One explanation could be that some degree of copolymerisation between the phenolic compounds and the pyrrole monomer unit took place. In a previous study conducted by this group, it was found that catechol was readily copolymerised with polypyrrole in the presence of ammonium persulphate and aluminium flake verified via XPS measurements [30]. 13.3.2

Density tests for reactions 1–6

The density tests provided information regarding the nature of the product. From Fig. 13.11, it is possible to observe the production of any free product for the

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13.5 SEM micrographs of ground product from reaction 2, at original magnifications of x 550 (top) and x 5,500 (bottom)

reactions 1 through 6. For example, any free product, product that is not attached to aluminium flake, will float in the liquid. The aluminium flakes and any adhered product will sink in the liquid due to higher density. It should be noted that the black line on the top of the liquid is the reflection of the product in the meniscus of the liquid. Reaction 1 product was tightly adhered to the flake as the entire product settled to the bottom of the liquid. Reaction 1x product has relatively good adherence. However, there is a higher occurrence of the free particulates in the liquid when compared to the product from reaction 1. Reaction 2 product appeared to form larger particulates that became adhered to the glass of the container, therefore, results were inconclusive. The product from reaction 3 is somewhat of an anomaly. It is extremely voluminous when compared with the other product as noticed during the filtering step of the synthesis. Indeed, reaction 3 product does display rather interesting structures that have a tissue-like appearance as observed via SEM. The density test would seem to indicate that, indeed, the product is quite different from the other products as shown by the low-density, floating behaviour observed during the test.

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13.6 SEM micrographs of ground product from reaction 3, at original magnifications of x 550 (top) and x 5,000 (bottom)

It is interesting that the micrographs indicating uneven deposition and the appearance of rather poorly adherent product did not directly correspond to the findings of free product during the density experiment. Reactions 1, 1x, 4, 4x, 5, and 6 display similar density results with reaction 5 possessing the highest amount of free product as determined visually. When comparing the micrographs, there would appear to be a large difference in free product in the reactions. Perhaps some of the small particulates observed in the SEM micrographs are not free product as thought, but instead small aluminium particles that resulted from the breakup of some of the larger flakes. This would effectively create small particles that could subsequently undergo deposition. Additionally, the small particles could be actually quite well attached to the surface and are not as easily removed as initially thought. 13.3.3

X-ray photoelectron spectroscopy of reaction 1

The XPS survey scan of the as-received Stapa Aloxal PM 2010 can be seen in Fig. 13.12. The majority of the material present in the as-received flake is revealed to

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13.7 SEM micrographs of ground product from reaction 4, at original magnifications of x 550 (top) and x 5,000 (bottom)

be comprised of aluminium, oxygen and carbon. It is likely that the large amount of oxygen in the sample is due to the Al2O3 on the surface of the flake which forms spontaneously when aluminium is exposed to the atmosphere. The sizeable quantity of carbon present can be attributed to the 1-methoxy-2-propanol (MP) on the surface of the flakes which is used as a surface treatment by the manufacturer to enable the flakes to become easily dispersed in water-borne coatings. The composition of the aluminium flake pigment according to the Material Safety Data Sheet (MSDS) supplied by the manufacturer of the aluminium flake is approximately 26% aluminium oxide, 39% aluminium, and 35% 1-methoxy-2-propanol by weight, respectively. The larger amount of oxygen and smaller amount of aluminium in the sample versus what is stated in the MSDS could be due to the oxidation of the aluminium flake after manufacturing. In addition, the difference between the amounts of carbon observed using XPS and the values given in the MSDS could be due to the volatility of the 1-methoxy-2-propanol. The core level carbon 1s (C1s) nitrogen 1s (N1s), and aluminium 2p (Al2p) XPS spectra of the polypyrrole/Al flake hybrid prepared in the presence of catechol, reaction 1, are shown in Fig. 13.13. A standard peak shape analysis with Gaussian fitting

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13.8 SEM micrographs of ground product from reaction 4x, at original magnifications of x 1,000 (top) and x 5,000 (bottom)

was used to analyse the core level spectra. Analysis revealed that the C1s main peak is split into three peaks as is shown in Fig. 13.13a. The peak at 284.98 eV is associated with the pyrrole Ż carbons, while the peak at 286.85 eV is associated with the pyrrole ź carbons. The energy gap of 1.87 eV between the two peaks is consistent with, but slightly larger than what has been previously reported for similar samples [31–33]. ‘Disorder effects’ such as inter-chain links, side chains or chain terminations can be allocated to the asymmetry of the C1s spectrum [31,32]. This is not unexpected when a strong oxidant such as ammonium persulphate is used as free radicals formed on β carbons in the pyrrole rings would lead to cross-linking between chains via the β carbons. As a result of these disorder effects, the electron density normally attributed to the hydrogen atom binding with Ż carbon of the pyrrole ring is mainly attributed to the Ż carbon itself. The result is a peak width which is wider in comparison to that of an ź carbon, as a result of the higher binding energy induced by the disorder effects. This can be seen in the peak at 288.97 eV in Fig. 13.13a. The N(1s) spectrum shown in Fig. 13.13b is characteristic of oxidised polypyrrole complexes that contain only amine-like and N+ nitrogen atoms with peaks at 399.41 eV and 401.31 eV, respectively. The electrostatic effect of the nearest counter

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13.9 SEM micrographs of ground product from reaction 5, at original magnifications of x 2,000 (top) and x 5,000 (bottom)

ion to the nitrogen can be assigned to the shoulder on the side of the peak that is higher in energy [31,34]. The nitrogen in the polypyrrole/aluminium hybrid flakes is composed of approximately 31% N+ which is similar to levels reported by other research groups [31,34]. There is not a peak present that can be assigned to an imine nitrogen. This indicates that the polypyrrole did not undergo a deprotonation step and an anion is still acting as a counter ion to the polymer [34]. Another possible explanation for the lack of an imine peak suggested by Lim et al., would be the formation of complexes between the nitrogen on the pyrrole rings and the aluminium along with other atoms such as oxygen converting the nitrogen to an amine-like state. This would cause the disruption of the Ɖ-electron conjugation resulting in a N-Ɖ-AlOz peak. The peak associated with N-Ɖ-AlOz is observed at 74.81 eV in Fig. 13.13c and could also be a contributing factor responsible for the large amount of oxygen observed [34]. The elemental composition of the polypyrrole/Al flake hybrid pigment prepared in reaction 1 can be seen in the survey spectra presented in Fig. 13.14. The reported ratio between the sulphur and nitrogen [S]/[N], is approximately 60%. Since Millipore

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13.10 SEM micrographs of ground product from reaction 6, at original magnifications of x 550 (top) and x 5,000 (bottom)

water was used in the synthesis of the PPy/Al hybrid and no ionic compounds were added to the synthesis other than the ammonium persulphate, it can safely be assumed that the counter ions present in the polypyrrole are the sulphate ions from the ammonium persulphate. If this assumption is indeed accurate, the [S]/[N] ratio implies that sulphate is the dopant anion existing in the ratio of three sulphate anions for every five pyrrole rings [31,35–37]. This hypothesis is supported by the S(2p) peak at 169.42 eV which can be attributed to the oxidised SO3– anion [38]. However, the oxygen to sulphur ratio is 11, much greater than the ratio of three, which is dictated by the structure of the APS. There are several explanations for this discrepancy. One possible explanation originates from the carbon/nitrogen [C]/[N] ratio which is approximately 7:1. This suggests that a copolymer could possibly be forming that contains both the polypyrrole and catechol with a ratio of one catechol ring for every two pyrrole rings in the polymer chain. The hydroxyl groups present on the catechol would help to explain the increased [O]/[S] ratio. Still another explanation originates from the aluminium flakes themselves. Since it is almost a certainty that the aluminium flakes were not totally covered by the polypyrrole and the flake was exposed to the ammonium persulphate before the addition of the catechol and pyrrole, it is

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13.11 Density test results for the ground product from reactions 1 through 6

13.12 XPS plot of the as-received flake

conceivable that a layer of Al2O3 covering the flakes would also be responsible for the large amount of oxygen present in the hybrid pigment. The fluorine peak is most likely an impurity in the aluminium flake as this peak was present in the XPS spectra of the pure aluminium flake.

Polypyrrole/aluminium flake hybrid pigments for corrosion inhibition

13.13 The core level spectra of (a) C 1s, (b) N 1s, and (c) Al 2p spectra of the PPy/Al flake hybrid synthesised in the presence of catechol, reaction 1

254

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13.14 XPS plot of the PPy/Al flake hybrid flake from reaction 1

13.3.4

Conductive atomic force microscopy (AFM) of reaction 1

To help verify that the product deposited onto the flake was indeed conductive, conductive AFM measurements were taken for the flakes from reaction 1. As can be seen in Figs. 13.15 and 13.16, the aluminium flakes have the majority of their surfaces covered in a conductive layer. The average conductance of the deposited product was 1.6 S/cm. This conductivity is due to the polymeric film deposited on the surface of the flake, as the conductivity of the as-received flake was too low to be detected when examined via C-AFM due to the resistive nature of the oxide layer. 13.3.5

Electrochemical impedance spectroscopy (EIS) of reaction 1

EIS measurement of a coated substrate is typically used to determine the barrier properties of the coating as well as monitor the degradation process of the system. The use of EIS as a formulating tool in determining the critical pigment volume

13.15 A 10 μm (left) and 5 μm (right) conductive AFM image showing the conductivity of the product from reaction 1

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13.16 A 2.5 μm conductive AFM image showing the conductivity of the product from reaction 1

concentration (CPVC) of a coating was first suggested by Lobnig et al. in 2006 [39]. This technique involves detecting the point at which the low frequency impedance drops due to the presence of voids in the coating near the CPVC. While this would have a definite advantage over many other commonly used methods for determining CPVC, it would seem to be limited to nonconductive pigments. In the case of conducting pigments, the percolation threshold, the point at which the pigment particles start to form electrical connections throughout the coating would be detected by a drastic drop in impedance while the CPVC would be largely unnoticed. When formulating a primer for corrosion prevention using a conductive pigment such as zinc for the cathodic protection of steel, the percolation threshold is a critical point that must be exceeded if an electrical connection is to be made with the substrate. Because a top coat would be applied over the primer for barrier properties, one would be more concerned with far exceeding the percolation threshold and less concerned with porosity introduced by exceeding the CPVC. This concept is illustrated in Fig. 13.17. The measurements were taken from epoxy primers containing either non-conductive as-received aluminium flake or a conductive PPy/Al flake hybrid pigment from reaction 1. Before EIS measurements were taken, the samples were immersed in dilute Harrison’s solution for 3 h. Notice that for the conductive PPy/Al flake primer, there is a large drop in impedance between a PVC of 30% and 35% while this same drop is not observed for the as-received flake until a PVC of 45%. The drop observed at the lower PVC for the conductive pigment would be indicative of the percolation threshold while the impedance drop observed at the higher PVC for the non-conductive pigment would be more closely associated with the CPVC. Also it is important to note that at a PVC of 35%, the impedance of the PPY/Al flake hybrid primer drops below the impedance observed for bare aluminium 2024. This is indicative of some interaction between the aluminium and PPy at the interface between the aluminium substrate and the coating. 13.3.6

Corrosion assessment of reaction 1

As can be seen in Fig. 13.18, the open circuit potential for a panel coated with a PPy/Al flake hybrid pigment, synthesised in the presence of catechol, is more positive

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13.17 Graph of low frequency impedance at 0.01 Hz vs. PVC

13.18 Plot of open circuit potential vs. immersion time in B117 salt spray

than for both uncoated aluminium 2024 and aluminium 2024 coated with an as received flake primer. This would support the observations of Herrasti and Ocon as no attempt was made to include corrosion inhibiting counter ions into the polypyrrole. Therefore, the more positive open circuit potential is the most probable cause for corrosion protection afforded to the aluminium 2024 substrate [40]. The corrosion inhibiting behaviour expected from these measurements was confirmed by exposing coated aluminium 2024 panels at a PVC of 45% to B117 exposure. The results of this exposure can be seen in Fig. 13.19. While the as-received flake sample developed blisters and began to fail after only 100 h of exposure, the PPy/Al flake primer did not develop any blisters. Some corrosion product is present in the scribe

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13.19 Sample coated with the PPy/Al flake hybrid primer (left) and a sample coated with the as-received flake primer (right) after 1300 h of B117 exposure

but there is a minimal amount of creep along the scribes indicating the primer’s corrosion inhibiting nature. Figures 13.20 and 13.21 are Bode plots showing EIS measurements taken periodically throughout the exposure of the samples. It is interesting to note that while there is very little corrosion present on the PPy hybrid sample, the impedance of that coating is very low. This is not entirely unexpected given the conductive nature of the pigment added to the epoxy, but it highlights the importance of knowing the nature of the sample that is being measured when using electrochemical techniques. Given no other information about the samples, it might be assumed that the coating with the higher impedance is the better coating. This, however, would be an incorrect interpretation of the results. 13.4

Conclusions

Polypyrrole composites were prepared under various experimental conditions. It was determined through SEM that the deposited material depended greatly on the reactants used during the synthesis. Reactions 1, 1x, and 3 resulted in flakes that were joined together. Reactions 4 and 4x resulted in small particulates which preferentially deposited on the edges of the aluminium flake. Reaction 5 produced sporadic coverage of the flakes, with most of the flake being uncoated by the deposited product. Reaction 6 produced microstructures which will be examined further in future work. The density test determined that most of the product was attached to the flake for reactions 1, 1x, 2, 4, 4x, 5, and 6. Reaction 3 displayed unique results where all of the product floated in the liquid, which is most likely related to the tissue-like appearance of the product that was found during the SEM study.

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13.20 Bode plot showing the change in impedance of an as-received flake coating, at a PVC of 45%, as a result of B117 salt spray exposure

13.21 Bode plot showing the change in impedance of a PPy/Al flake hybrid coating, at a PVC of 45%, as a result of B117 salt spray exposure

For the XPS, C-AFM, open circuit potential (OCP), accelerated exposure (B117), and electrochemical impedance spectroscopy (EIS) experiments, the study focused on the product from reaction 1. It was found that the product was most likely a co-polymer formed between the pyrrole and catechol. The results from reaction 1 also seemed to show that the dopant was sulphate ions from the ammonium

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persulphate. The C-AFM detected a conductive surface with a conductivity of 1.6 S/cm. EIS was used to assist in the formulating of a conductive coating where it was determined that a 45% PVC would be sufficient for a low impedance coating of the polypyrrole composite flake. The OCP measurements detected a positive shift in the open circuit potentials for the aluminium alloy coated with the polypyrrole composite flake during the accelerated exposure in B117. Additionally, the samples exposed to B117 displayed very different properties when examined visually with the polypyrrole composite flakes outperforming the as-received flake system. The EIS results indicate that the as-received flake coating is formulated under its critical PVC as the coating initially displays relatively high impedance at low frequencies. This low-frequency impedance decreases over the duration of the experiment, indicating a decrease in the barrier protection provided by the coating systems. In comparison, the impedance spectra for the polypyrrole composite flake are relatively low throughout the exposure time. However, no visible signs of corrosion were observed during the study period. Future studies will probe the corrosion inhibition by the products of the other reactions: 1x, 2, 3, 4, 4x, 5, and 6 to determine which reaction conditions produce the composite flake with greatest corrosion inhibition. Acknowledgments The authors would like to thank the US Army Research Laboratory (Contract #: W911NF-04-2-0029) for sponsoring this research. References 1. H. Shirakawa, E. J. Louis, A. G. MacDiarmid, C. K. Chiang and A. H. Heeger, J. Chem. Soc. Chem. Commun. (1977), 578–580. 2. ‘The Nobel prize in chemistry, 2000: conductive polymers’, The Royal Swedish Academy of Sciences (2000), 1–15. 3. D. W. DeBerry, J. Electrochem. Soc., 132(5) (1985), 1022–1026. 4. T. P. McAndrew, S. A. Miller, A. G. Gilicinski and L. M. Robeson, Proc. Am. Chem. Soc. Polym. Mater. Sci. Eng., 74 (1996), 204–206. 5. J. O. Stoffer, S. P. Sitaram and T. J. O’Keefe, J. Coat. Technol., 69 (1997), 65–69. 6. I. V. Yagova, S. S. Ivanov, I. V. Bykov and V. V. Yagov, Prot. Met., 34 (1998), 132–136. 7. W.-K. Lu, S. Basak and R. L. Elsenbaumer, ‘Corrosion protection of metals using electrically conductive polymers’, in Handbook of Conductive Polymers, 881–920, ed. T. A. Skotheim, J. Reynolds, R. L. Elsenbaumer. Marcel Dekker, Inc., New York, NY, 1997. 8. D. E. Tallman, C. Vang, G. G. Wallace and G. P. Bierwagen, J. Electrochem. Soc., 149(3) (2002), C173–C179. 9. G. M. Spinks, A. J. Dominis, G. G. Wallace and D. E. Tallman, J. Solid State Electrochem., 6 (2002), 85–100. 10. G. Paliwoda-Porebska, M. Rohwerder, M. Stratmann, U. Rammelt, L. M. Duc and W. Plieth, J. Solid State Electrochem., 10 (2006), 730–736. 11. D. E. Tallman, G. Spinks, A. Dominis and G. G. Wallace, J. Solid State Electrochem, 6 (2002), 73–84. 12. J. Ruhe, T. Ezquerra and G. Wegner, Makromol. Chem., 10 (1989), 103–108. 13. A. O. Patil, Y. Ikenoue, F. Wudl and A. J. Heeger, J. Am. Chem. Soc., 109 (1987), 1858–1859. 14. A. A. Ashraf, F. Chen, C. O. Too and G. G. Wallace, Polymer, 37(13) (1996), 2811–2819.

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15. W. Prissanaroon, N. Brack, P. J. Pigram and J. Liesegang, Synth. Met., 142 (2004), 25–34. 16. A. Malinauskas, Polymer 42 (2001), 3957–3972. 17. J. I. Martins, T. C. Reis, M. Bazzaoui, E. A. Bazzaoui and L. Martins, Corros. Sci., 46 (2004), 2361–2381. 18. E. Armelin, R. Pla, F. Liesa, X. Ramis, J. I. Iribarren and C. Aleman, Corros. Sci., 50 (2008), 721–728. 19. E. Armelin, R. Oliver, F. Liesa, J. I. Iribarren, F. Estrany and C. Aleman, Prog. Org. Coat., 59 (2007), 46–52. 20. C. A. Amarnath, S. Palaniappan, P. Rannou and A. Pron, Thin Solid Films, 516 (2008), 2928–2933. 21. D. Aradilla, C. Ocampo, E. Armelin, C. Aleman, R. Oliver and F. Estrany, Mater. Corros., 58(11) (2007), 867–872. 22. S. Sathiyanarayanan, S. Muthukrishnan, G. Venkatachari and D. C. Trivedi, Prog. Org. Coat., 53 (2005), 297–301. 23. S. Sathiyanarayanan, S. Syed Azim and G. Venkatachari, Synth. Met., 157 (2007), 751–757. 24. S. Sathiyanarayanan, S. Syed Azim and G. Venkatachari, Electrochim. Acta, 53 (2008), 2087–2094. 25. S. Sathiyanarayanan, S. Syed Azim and G. Venkatachari, Prog. Org. Coat., 59 (2007), 291–296. 26. K. L. Levine, D. E. Tallman and G. P. Bierwagen, Aust. J. Chem., 58 (2005), 294–301. 27. T. Dai, X. Yang and Y. Lu, Nanotechnology, 17(12) (2006), 3028–3034. 28. X. Zhang, J. Zhang, Z. Liu and C. Robinson, Chem. Commun. R. Soc. Chem. (2004), 1852–1853. 29. G. Lu, C. Li and G. Shi, Polymer, 47 (2006), 1778–1784. 30. C. Vetter, X. Qi, A. Harper, S. Kasisomayajula and V. J. Gelling, in Proceedings of International Corrosion Conference, 6–10 October 2008, Las Vegas, NV. 31. J. Joo, J. K. Lee, S. Y. Lee, K. S. Lang, E. J. Oh and A. J. Epstein, Macromolecules, 33 (2000), 5131–5136. 32. P. Pfluger and G. B. Street, J. Chem. Phys., 80 (1984), 544–553. 33. L. Atanasoska, K. Naoi and W. H. Smyrl, Chem. Mater., 4 (1992), 988–994. 34. V. W. L. Lim, S. Li, E. T. Kang, K. G. Neoh and K. L. Tan, Synth. Met., 106 (1999), 1–11. 35. X. Zhang, J. Zhang, W. Song and Z. Liu, J. Phys. Chem. B, 110 (2005), 1158–1165. 36. W. Prissanaroon, N. Brack, P. J. Pigram and J. Liesegang, Synth. Met., 142 (2004), 25–34. 37. M. Hepel, Electrochim. Acta, 41(1) (1996), 63–76. 38. C. Malitesta, I. Losito, L. Sabbatini and P. G. Zambonin, J. Electron Spectrosc. Relat. Phenom., 97 (1998), 199–208. 39. R. E. Lobnig, W. Villalba, K. Goll, J. Vogelsang, I. Winkels, R. Schmidt, P. Zanger and J. Soetemann, Prog. Org. Coat., 55 (2006), 363–374. 40. P. Herrasti and P. Ocon, Appl. Surf. Sci., 172 (2001), 276–284.

14 Electrochemical behaviour of ZrO2 sol–gel films doped with corrosion inhibitors on AA2024 aluminium alloy F. Andreatta, L. Paussa, P. Aldighieri and L. Fedrizzi Department of Chemical Science and Technology, University of Udine, Via del Cotonificio 108, 33100 Udine, Italy [email protected]

14.1

Introduction

Aluminium alloys for aerospace applications are protected with a paint system consisting of a conversion layer, primer and top coat. Cr-based pre-treatments are extensively employed as conversion coatings because they provide very good adhesion for primer and top coat associated with good barrier properties [1]. Moreover, chromate conversion coatings exhibit self-healing ability [1]. However, hexavalent chromium-containing compounds are regarded as toxic and not environmental friendly [1–4]. Therefore, use of chromate conversion coatings is restricted and alternative pre-treatments should be employed for aluminium alloys. Coatings obtained with the sol–gel technique are a possible replacement for chromate conversion coatings. This technique enables deposition of mixed metal oxides and metal oxide–organic composites on different metal substrates [1,4–7]. The thickness of coatings deposited with sol–gel technology is generally in the micrometre range (1–20 μm). Corrosion protection of different types of sol–gel coatings has been extensively investigated in the past [1]. Oxide layers deposited by dip-coating and spin-coating from sol–gel systems improve the corrosion resistance of aluminium alloys [1,8]. Silane-based sol–gel coatings are reported to improve resistance of aluminium to general and localised corrosion [9,10]. Moreover, mixed SiO2 and ZrO2 oxide layers deposited on AA2024 with the sol–gel technique increased resistance to localised attack [11]. In addition, it was shown that Zr-based sol–gel coatings can be used as barrier layers in paint systems for AA2024 [12]. Sol–gel technology presents the possibility of introducing corrosion inhibitors in the oxide layer. Cr-free organic and inorganic inhibitors have been incorporated in different sol–gel coatings to improve corrosion resistance and provide self-healing ability. Triazole and thiazole compounds have been investigated by several authors [13–15]. Triazole compounds promote strong cathodic inhibition because they tend to be adsorbed on Cu-rich cathodic phases [13]. Moreover, they might also affect anodic processes on the alloy surface [15]. This behaviour is explained by the ability of triazole and thiazole compounds rapidly to form a protective film on the alloy surface. The introduction of active nanoparticles such as ZrO2 in sol–gel coatings is another interesting approach to improve corrosion resistance of AA2024 [16]. Ce-based conversion coatings exhibited promising self-healing ability in different studies [17]. Despite the fact that the self-healing ability of Ce-based pre-treatments is not as good as chromate conversion coatings, inorganic Ce compounds such as 262

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cerium nitrate might be considered for application as corrosion inhibitors in sol–gel systems. The introduction of cerium nitrate in Zr-based sol–gel coatings might be difficult because it might affect reactions involved in the deposition of the sol–gel coating [18]. This might affect the barrier properties of sol–gel layers doped with cerium nitrate [16]. Nevertheless, cerium nitrate was shown to have a positive effect on long-term corrosion resistance of Si-based sol–gel coatings [16]. In the case of Zr-based sol–gel coatings, it was not possible to determine a clear inhibition effect from the incorporation of cerium nitrate [18,19]. The development of non-chromated coating systems for aerospace application has been previously targeted with the attempt to replace chromate conversion coating with a thin sol–gel layer (typically 50–200 nm) [20]. This approach is new because sol–gel coating systems are usually thick (in the micrometre range) and are generally regarded as complete coating systems. The superior adhesion behaviour of thin Zr- and Si-based sol–gel films relative to chromate conversion coatings is the main advantage of developing thin coatings. Barrier properties and long-term corrosion resistance need to be further improved to substitute chromate conversion coatings with thin sol–gel films. Moreover, introduction of inhibitors in thin sol–gel films is still a critical aspect [20]. In previous work, our research group followed the strategy of developing thin sol–gel coatings for different aluminium alloys [21,22]. Thin ZrO2 amorphous films (100–200 nm) were deposited on AA1050 with the dip-coating technique [21]. These coatings promote adhesion of organic coatings to the ZrO2 pretreatment. Moreover, the electrochemical behaviour of AA1050 coated with ZrO2 films is strongly dependent on deposition parameters, which affect film quality. ZrO2pre-treated AA6060 exhibits barrier properties similar to chromate conversion-coated AA6060 when a continuous layer was deposited on the surface through successive dipping steps in the sol–gel solution [22]. ZrO2 pre-treatments on AA6060 do not exhibit self-healing ability like chromate conversion-coated AA6060. However, a limited recovery of the barrier properties was observed for ZrO2 pre-treatments with a small number of defects. This was attributed to the formation of corrosion products that might plug the defects in the layer. The work presented in this paper was carried out in the framework of the MULTIPROTECT project (funded by the European Community as contract N° NMP3-CT-2005-011783). It deals with the deposition of thin (100–200 nm) ZrO2 sol–gel pre-treatments on AA2024 and evaluation of their barrier properties. Due to the limited thickness of the layers deposited in our work (70–180 nm), it is very important to evaluate the ability of ZrO2 sol–gel pre-treatments to provide a good barrier against corrosion. Moreover, the introduction of corrosion inhibitors (2-mercaptobenzothiazole and cerium nitrate) is targeted to improve corrosion resistance. 14.2

Experimental

The substrate for the deposition of ZrO2 films was AA2024 aluminium alloy. The alloy composition is given in Table 14.1. Table 14.1 Composition of AA2024 aluminium alloy (wt.%) Al Bal.

Cu

Mg

Si

Fe

Mn

Zr

Ti

Cr

3.8–4.9

1.2–1.8

<0.5

<0.5

0.3–0.9

<0.25

<0.15

<0.1

Electrochemical behaviour of doped ZrO2 sol–gel films

264

Before deposition of the sol–gel layers, the sample surface underwent preparation consisting of alkaline cleaning, alkaline etching and acid etching. The sol–gel solutions were 0.1 M and 0.6 M Zr(OBun)4 in anhydrous n-butanol. The deposition of the film was performed by the dipping sol–gel technique. The samples were withdrawn at a controlled rate of 1 mm/s. The dipping step was repeated 2, 3 or 4 times to deposit overlapping layers on AA2024 samples. After each dip in the sol–gel solution, the samples were subjected to thermal treatment. This was performed at 120ºC for 4 min. Thermal treatment was followed by drying at room temperature. Two corrosion inhibitors were selected for doping ZrO2 pre-treatments deposited with sol–gel technology: 2-mercaptobenzothiazole and cerium nitrate. The introduction of inhibitors in sol–gel pre-treatments was carried out following two approaches. In the first approach, the inhibitor was added to the sol–gel solution leading to the formation of a sol–gel film doped with the inhibitor. In the second, the inhibitors were added with a final dip in a solution containing the inhibitor on a multilayer structure consisting of a first layer of ZrO2 obtained with 0.1 M sol–gel solution and a second layer deposited with 0.6 M solution. The inhibitors were alternatively applied by spraying in the second approach. The morphology of ZrO2 sol–gel layers was investigated using a scanning electron microscope (SEM) equipped with energy dispersive X-ray spectroscopy (EDXS). Glow discharge optical emission spectroscopy (GDOES) depth profiles were recorded with a Jobin Yvon GD profiler. The source parameters were set at 40 W applied power and 650 Pa Argon pressure. Quantitative profiles were obtained with a calibration procedure using certified reference materials. The electrochemical behaviour of samples coated with ZrO2 was investigated by potentiodynamic polarisation measurements and electrochemical impedance spectroscopy (EIS) in 0.05 M NaCl solution. Potentiodynamic polarisation measurements were recorded using a Pt counter electrode and Ag/AgCl reference electrode. The area of the working electrode was 3.6 cm2. The scan rate was 0.2 mV/s. Impedance measurements were carried out at open circuit potential with an AC voltage amplitude of 10 mV and frequency range from 10 mHz to 100 kHz. The measurements were performed using an AUTOLAB PG-STAT 12 potentiostat equipped with a Frequency Response Analyser module. In order to compare the barrier properties of samples coated with ZrO2, potentiodynamic polarisation scans and impedance measurements were performed on chromate conversion-coated AA2024. 14.3

Results and discussion

Figure 14.1A shows an SEM micrograph of the surface of as-received AA2024. The rolling direction is clearly visible in the micrograph. Moreover, second phase particles appearing darker than the matrix can be seen in the same micrograph. These particles are Al–Cu–Fe–Mn–Si intermetallics, which are typical constituent particles of AA2024 [23]. It is also possible that a number of dark spots are associated with holes generated by removal of intermetallics during hot- and cold-rolling of the alloy. Before deposition of the sol–gel film, the samples underwent surface preparation consisting of alkaline degreasing, alkaline etching and acid etching. The aim of surface preparation was to remove the intermetallics and the defective oxide layer resulting from cold-rolling. Figure 14.1B shows the surface of AA2024 after surface preparation. Alkaline etching leads to the formation of rather large cavities on the

265

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14.1 SEM micrographs of as-received AA2024 (A) and of alloy surface after surface preparation

alloy surface. This is probably related to non-uniform attack due to the inhomogeneous microstructure of AA2024 alloy. Moreover, it is likely that cavities are generated by removal of intermetallics during alkaline etching. Figure 14.2A is an SEM micrograph of AA2024 coated with a ZrO2 film deposited with the sol–gel technique (4 layers). The ZrO2 film can not be clearly distinguished in the micrograph since the deposited layer is thin (in the range of 100–200 nm). However, the SEM micrograph at higher magnification in Fig. 14.2B reveals the existence of a ZrO2 film, which appears uniform and well adherent on the substrate. This was further confirmed by TEM observation of cross-sections of samples coated with a ZrO2 film. This indicates that AA2024 is uniformly covered by deposition of a ZrO2 film with the sol–gel technique. Figure 14.3 displays a semi-quantitative GDOES composition profile for a sample coated with 2 layers of ZrO2 deposited with the sol–gel technique. The Zr and O signals reveal the existence of a layer of ZrO2 on the substrate. The interface between

Electrochemical behaviour of doped ZrO2 sol–gel films

266

14.2 SEM micrographs of AA2024 samples pre-treated with ZrO2 layers deposited with the sol–gel technique

the sol–gel film and the aluminium substrate is associated with a decrease in Zr and O signals and an increase in Al signal. This is located at about 150 nm indicating that the sol–gel technique can be employed for the deposition of thin ZrO2 layers on AA2024. This result is in line with SEM micrographs reported in Fig. 14.2. Moreover, the transition of Al, Zr and O across the interface is rather broad. This is probably related to the surface roughness of coated samples (Fig. 14.2A and B), which might affect GDOES measurements. Moreover, the GDOES profile in Fig. 14.3 exhibits a C signal that is relatively high in the ZrO2 film and progressively decreases approaching the substrate. The existence of a C signal in the sol–gel layer indicates that a fraction of the metal-organic precursor remains in the film. In addition, contamination by carbon is also possible on the sample surface. ZrO2 films become thicker with increasing numbers of dips in the sol–gel solution, reaching more than 200 nm after 4 dips in the solution. However, GDOES semi-quantitative profiles reveal that the increase in film thickness is limited for

267

Self-healing properties of new surface treatments

14.3 GDOES semi-quantitative composition profile for ZrO2 pre-treatment on AA2024

depositions after the first one. The thickness of these sol–gel ZrO2 films is in the same range as that of conventional chromate conversion coatings on AA2024, which are generally in the range of 100–300 nm. From this point of view, these sol–gel layers might be an option for the replacement of Cr-based conversion coatings, as also indicated by our previous work on AA1050 and AA6060 [21,22]. Figure 14.4 shows typical defects that might be observed in thin ZrO2 pretreatments deposited with the sol–gel technique. Figure 14.4A shows a discontinuity in the sol–gel layer with size of about 5 μm. The size of this discontinuity is similar to features observed on the substrate after surface preparation (Fig. 14.1B). It is possible that the sol–gel solution could not penetrate inside smaller cavities leading to a defective film structure at these locations. Figure 14.4B reveals the existence of cracks in the ZrO2 film. These cracks tend to be located at the bottom of relatively large cavities because a thicker film is formed within the cavities than outside, as confirmed by EDXS maps. Indeed, there is a larger amount of sol–gel solution at the bottom of large cavities than in the surrounding regions when samples are located in the horizontal position in the oven for thermal treatment of the sol–gel layer. Surface preparation before film deposition with the sol–gel technique is very important in order to obtain a uniform ZrO2 layer on the substrate. The procedure for surface preparation of the substrate needed to be optimised to reduce the number of defects in the sol–gel layer. Indeed, film defects might impair barrier properties leading to poor corrosion resistance. In order to assess the barrier properties of ZrO2 pre-treatments, potentiodynamic polarisation scans were initially carried out for undoped ZrO2 layers on AA2024. Figure 14.5 displays potentiodynamic polarisation curves determined in 0.05 M NaCl solution on AA2024 coated with 2 and 4 layers of ZrO2 and on bare substrate. The coated samples show a marked reduction in

Electrochemical behaviour of doped ZrO2 sol–gel films

268

14.4 SEM micrographs showing typical defects on AA2024 samples pre-treated with ZrO2 layers deposited with the sol–gel technique

cathodic and anodic current density with respect to the bare substrate. This can be explained by the reduction in the area available for cathodic and anodic reactions due to the ZrO2 film. Extrapolation of corrosion current density from Tafel slopes gives 1·10–6 A cm–2 for AA2024 coated with 2 layers of ZrO2 and 2·10–7 A cm–2 for pretreatment with 4 layers. In contrast, the corrosion current density is 1·10–5 A cm–2 for the bare substrate. The strong reduction in corrosion current density of coated samples with respect to bare AA2024 indicates that deposition of the ZrO2 pretreatment with the sol–gel technique leads to a marked improvement in the corrosion resistance of AA2024. This is due to the barrier effect of the deposited film. This is evident from the value of anodic current density at –0.5 V vs Ag/AgCl (immediately after the corrosion potential), which is three orders of magnitude lower for pretreatment with 4 layers of ZrO2 (5.5·10–7 A cm–2) than for bare substrate (5.5· 10–4 A cm–2). The very low anodic current density for the pre-treated sample is an

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Self-healing properties of new surface treatments

14.5 Potentiodynamic polarisation curves in 0.05 M NaCl solution for AA2024 samples pre-treated with ZrO2 layers deposited with the sol–gel technique

indication of the strong reduction in the corrosion of aluminium due to the presence of the ZrO2 film. Nevertheless, the potentiodynamic polarisation curve for a sample coated with 4 layers of ZrO2 does not exhibit a clear passive behaviour in the anodic branch. This behaviour might be related to small number of defects present in the ZrO2 layer (Fig. 14.4). Film defects might be preferential sites for localised attack. This is confirmed by the SEM micrograph reported in Fig. 14.6 showing the surface

14.6 SEM micrograph of an AA2024 sample coated with ZrO2 deposited with the sol–gel technique after potentiodynamic polarisation in 0.05 M NaCl solution

Electrochemical behaviour of doped ZrO2 sol–gel films

270

of coated AA2024 after potentiodynamic polarisation in 0.05 M NaCl. The micrograph exhibits a relatively large pit with a size of about 100 μm, while the surrounding area does not appear to have been attacked. In addition, the ZrO2 film displays good adhesion to the substrate after potentiodynamic polarisation. The corrosion attack most likely initiates at discontinuities similar to those reported in Fig. 14.4. Moreover, the attack is localised at film defects leading to pitting of the surface. Since defects in the ZrO2 film deposited with the sol–gel technique have a negative effect on the electrochemical behaviour, the introduction of corrosion inhibitors in the pre-treatment was considered as a way to improve corrosion resistance. Figure 14.7A shows an SEM micrograph of the surface of AA2024 coated with ZrO2 pre-treatment doped with 2-mercaptobenzothiazole. The inhibitor was introduced into the sol–gel solution before deposition by dipping of the ZrO2 layer. The sample

14.7 SEM micrographs of AA2024 coated with ZrO2 layers doped with 2mercaptobenzothiazole: (A) inhibitor introduced into the sol–gel solution before deposition with the sol–gel technique and (B) multilayer pre-treatment structured with a first ZrO2 layer obtained with 0.1 M precursor concentration, a second layer deposited with 0.6 M concentration and final addition of inhibitor

271

Self-healing properties of new surface treatments

in Fig. 14.7A is coated with 2 layers of ZrO2 doped with inhibitor. The film morphology is very similar to that shown by the undoped ZrO2 pre-treatment (Fig. 14.2A). The alloy surface appears uniformly coated revealing cavities produced during pickling. As already considered above, these film defects might be related to the alloy morphology obtained after surface preparation. However, film discontinuities might also be generated by inhibitor addition to the sol–gel bath leading to a viscosity increase. This renders film cracking more probable relative to thin films obtained from a precursor solution with low viscosity. The SEM micrograph in Fig. 14.7B shows the surface of a ZrO2 pre-treatment consisting of two overlapping layers of ZrO2. The first layer was obtained with 0.1 M precursor concentration in the sol–gel solution, as for the sample in Fig. 14.7A. The second was produced by dipping in 0.6 M sol–gel solution. The inhibitor was added to this multilayer system by a final dip in a solution containing 2-mercaptobenzothiazole or by spraying the same solution on the pre-treatment. In addition, samples doped with cerium nitrate were also produced by the same approach. The morphology of pre-treatments doped with cerium nitrate is similar to that shown in Fig. 14.7B. It can be seen that the second layer of pre-treatment (0.6 M precursor concentration) is discontinuous and porous. It can be expected that the first layer is similar to that visible in Fig. 14.2A since it was obtained according to the same procedure. The first layer should provide a barrier for corrosion protection, while the second layer should enhance adhesion of the top coat and provide a microstructure that favours inhibitor storage in the ZrO2 film. Figure 14.8 shows a GDOES semi-quantitative composition profile for a sample pre-treated with ZrO2 film doped with 2-mercaptobenzothiazole. The pre-treatment consisted of two overlapping layers of ZrO2 containing inhibitor, as for the sample

14.8 GDOES semi-quantitative composition profile for ZrO2 pre-treatment containing 2-mercaptobenzothiazole

Electrochemical behaviour of doped ZrO2 sol–gel films

272

shown in Fig. 14.7A. It was not possible to carry out GDOES profiles on the multilayer pre-treatment shown in Fig. 14.7B because surface roughness did not enable uniform sputtering disturbing signal acquisition. The semi-quantitative composition profile in Fig. 14.8 shows similar trends for Zr, O and Al signals to the profile for undoped pre-treatment. Film thickness is in the same range (150–200 nm) for doped and undoped pre-treatments. In general, the thickness of multilayer pre-treatment as in Fig. 14.7B is slightly greater than for the film in Fig. 14.7A. The S signal shows surface enrichment with a lower rather uniform concentration in the ZrO2 layer. This can be attributed to the incorporation of 2-mercaptobenzothiazole in the pre-treatment and is supported by the observation that the S signal is below detection limits for undoped pre-treatments (Fig. 14.3). It was only possible to obtain qualitative profiles for pre-treatments with a multilayer structure consisting of a layer obtained with 0.1 M precursor concentration, another obtained with 0.6 M concentration and a final addition of inhibitor (Fig. 14.7B). These profiles show the presence of S on the surface of the ZrO2 layer for samples doped with 2-mercaptobenzothiazole. Similarly, surface enrichment of Ce is detected for pre-treatments with additions of cerium nitrate. Figure 14.9 presents potentiodynamic polarisation curves for AA2024 coated with ZrO2 films containing 2-mercaptobenzothiazole. The inhibitor was added to the system by spraying on to the ZrO2 film deposited with the sol–gel technique. In addition, the figure displays the potentiodynamic polarisation curve for an undoped ZrO2 pre-treatment (2 layer) and for the bare AA2024 substrate. The curve for the ZrO2 pre-treatment doped with 2-mercaptobenzothiazole (applied by spraying) shows a passive range (about 200 mV) and breakdown at –0.37 V vs Ag/AgCl. Moreover,

14.9 Potentiodynamic polarisation curves in 0.05 M NaCl solution for AA2024 samples pre-treated with ZrO2 layers deposited with the sol–gel technique with and without 2-mercapotobenzothiazole

273

Self-healing properties of new surface treatments

it displays a marked reduction in cathodic and anodic current density with respect to the undoped pre-treatment and bare AA2024. The corrosion current density is 2·10–8 A cm–2 for the doped system. This is about two orders of magnitude lower than the corrosion current density for the undoped ZrO2 pre-treatment and about three orders of magnitude lower than that for bare substrate. The existence of a passive range associated with low anodic current density might be related to the deposition of a homogeneous barrier film (the layer deposited with 0.1 M precursor concentration). It is likely that this barrier layer is defect-free, in accordance with the low corrosion current density detected for this sample. The low cathodic current density for the same protection system is probably related to the addition of 2-mercaptobenzothiazole, which is known to be an efficient inhibitor of cathodic reactions. The combination of an efficient barrier layer and inhibition due to 2-mercaptobenzothiazole leads to good corrosion protection of the substrate during potentiodynamic polarisation. The results shown in Fig. 14.9 indicate that discontinuities in the ZrO2 pre-treatments deposited with the sol–gel technique are a critical factor concerning corrosion protection. Moreover, this has a negative effect on the reproducibility of potentiodynamic polarisation curves. This is consistent with the morphology observed for ZrO2 films. The electrochemical behaviour of pre-treatments obtained with direct addition of inhibitor in sol–gel solution (morphology shown in Fig. 14.7A) is similar or even worse than that for the undoped system. It is likely that the addition of 2mercaptobenzothiazole to the sol–gel solution negatively affects hydrolysis and condensation reactions involved in the deposition of the sol–gel layer, affecting film continuity. This might be reflected in the electrochemical behaviour. Figure 14.10 shows the potentiodynamic polarisation curve for a ZrO2 pretreatment doped with cerium nitrate by spraying. Again, the figure shows the curves

14.10 Potentiodynamic polarisation curves in 0.05 M NaCl solution for AA2024 samples pre-treated with ZrO2 layers deposited with the sol–gel technique with and without cerium nitrate

Electrochemical behaviour of doped ZrO2 sol–gel films

274

for pre-treatment without the inhibitor and for bare substrate. The cathodic and anodic current density is similar for pre-treatment with and without inhibitor. The same behaviour is observed for the corrosion current density. Moreover, there is a shift of the corrosion potential in the anodic direction for the pre-treatment with inhibitor suggesting a higher resistance with respect to passivity breakdown. Indeed, this potential might be considered as a breakdown potential in this case. This aspect needs to be investigated further. Electrochemical impedance measurements were carried out on ZrO2 pretreatments to evaluate corrosion protection for long immersion times in 0.05 M NaCl solution. Figure 14.11 displays the modulus of electrochemical impedance at 0.01 Hz versus immersion time for ZrO2 pre-treatment on AA2024 (2 layers and 4 layers). In addition, this figure shows the impedance modulus for a reference chromate conversion coating and for bare substrate. The impedance modulus measured at 0.01 Hz immediately after immersion in the electrolyte (0 h) is in the order of 106 ohm cm2 for the sample coated with 4 layers of ZrO2 and 3.7 105 ohm cm2 for that with 2 layers. It is 4.4 103 ohm cm2 for chromate conversion-coated AA2024 and 4.6.104 ohm cm2 for bare AA2024. Samples coated with ZrO2 film exhibit a decrease in impedance modulus during the first 8 h in solution. In contrast, the impedance modulus of chromate conversion-coated alloy progressively increases during the measurement. The impedance modulus is in the same range for the sample coated with ZrO2 (4 layers) and the chromate conversion-coated sample in the time interval between 24 h and 72 h. The impedance modulus is 2.9 105 ohm cm2 for chromate conversioncoated AA2024 after 192 h in solution, while it is about one order of magnitude lower for samples coated with ZrO2.

14.11 Electrochemical impedance modulus at 0.01 Hz as a function of immersion time in 0.05 M NaCl solution for AA2024 samples pre-treated with ZrO2 layers deposited with the sol–gel technique, chromate conversion-coated AA2024 and bare AA2024

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Self-healing properties of new surface treatments

The trend of impedance modulus as a function of immersion time shows that the barrier properties of ZrO2 film are higher than those of the chromate conversion coating for relatively short immersion times (about 24 h) in NaCl solution. As already seen from potentiodynamic polarisation curves, the sample coated with 4 layers exhibits a better barrier than that with 2 layers due to the existence of a thicker layer on the former sample. The sample with 4 layers of ZrO2 displays similar impedance values to the chromate conversion-coated sample between 24 h and 72 h immersion times indicating that the barrier properties remain rather good. In contrast, the chromate conversion-coated sample displays a higher impedance modulus than other samples for long immersion times. This is most likely due to self-healing that is well known for this type of coating [1]. This behaviour is not observed for samples coated with ZrO2 film. Figure 14.12 reports the electrochemical impedance modulus at 0.01 Hz versus immersion time for ZrO2 pre-treatments doped with 2-mercaptobenzothiazole. Chromate conversion coating and bare substrate are also reported in the figure. The impedance modulus measured at 0.01 Hz immediately after immersion in the electrolyte (0 h) is approximately 107 ohm cm2 for the sample coated with ZrO2 containing 2-mercaptobenzothiazole (indicated by number 2 in the figure). This sample is structured with two layers of ZrO2 oxide and the inhibitor was deposited directly on the pre-treated surface. The inner oxide layer was obtained with 0.1 M precursor concentration and the outer with 0.6 M, as described in the legend of Fig. 14.7B. The impedance modulus of this sample exhibits a rather rapid decrease over 48 h in the electrolyte. However, the modulus of the sample with ZrO2 pre-treatment doped with

14.12 Electrochemical impedance modulus at 0.01 Hz as a function of immersion time in 0.05 M NaCl solution for AA2024 samples pre-treated with ZrO2 layers deposited with the sol–gel technique with and without 2-mercaptobenzothiazole and chromate conversion-coated AA2024

Electrochemical behaviour of doped ZrO2 sol–gel films

276

2-mercaptobenzothiazole (sample 2) is significantly higher than that measured for the chromate conversion coating. Moreover, this is higher than the modulus of undoped ZrO2 film over 24 h in the electrolyte. The impedance modulus at 0.01 Hz is 5·105 ohm cm2 immediately after immersion (0 h) for the sample with inhibitor added to the sol–gel solution before pre-treatment application (indicated by number 1 in the figure). The impedance modulus for this sample shows a rapid decrease over 24 h in solution and becomes stable for longer immersion times. However, the impedance modulus of this sample is higher than that of the chromate conversion coating over 10 h in the electrolyte. The trend of the impedance modulus seen in Fig. 14.12 for the sample with ZrO2 pre-treatment doped with mercaptobenzothiazole shows a clear effect of the inhibitor on corrosion resistance only over a relatively short immersion time in the electrolyte (about 24 h). The inhibition is most likely associated with the deactivation of cathodic reactions, which is proven for this type of inhibitor [13]. This effect tends to disappear for longer immersion times. This might be related to complete release of the inhibitor over the first few hours in the electrolyte. It is believed that the release of inhibitor might be limited by the deposition of a paint onto the ZrO2 pretreatment. Indeed, the paint should favour retention of released inhibitor on the surface of the ZrO2 film, extending the positive effect of 2-mercaptobenzothiazole on electrochemical behaviour over longer times. Moreover, the samples containing 2-mercaptobenzothiazole do not exhibit signs of repassivation (self-healing ability) in contrast to the behaviour of the chromate conversion coating. Nevertheless, the long-term corrosion resistance of ZrO2 pre-treatment (with and without inhibitor) can be considered rather good taking into account that the film thickness is only in the range 100–200 nm. 14.4

Conclusions

Thin ZrO2 pre-treatments (100–200 nm) were applied by sol–gel technology to AA2024. The pre-treatments were doped with 2-mercaptobenzothiazole and cerium nitrate. Film morphology and composition have been investigated by SEM-EDXS and GDOES. Electrochemical behaviour has been investigated by potentiodynamic polarisation scans and electrochemical impedance spectroscopy to assess the barrier properties and long-term corrosion protection of the pre-treatments. •

•

The ZrO2 pre-treatment uniformly covers the substrate as confirmed by SEM micrographs and GDOES semi-quantitative composition profiles. Film morphology is strongly affected by surface preparation (alkaline cleaning, alkaline etch and acid etch) leading to the formation of rather large cavities due to inhomogeneous attack of the AA2024 substrate. This has a negative effect on the quality of pre-treatment, which might exhibit defects such as discontinuities or cracks. However, the number of defects in the ZrO2 pre-treatment can be minimised by modifying the deposition parameters. ZrO2 pre-treatments applied with the sol–gel technique improve the corrosion resistance of AA2024. This can be attributed to the barrier effect of the deposited pre-treatments. However, undoped pre-treatments do not show a passive range in potentiodynamic polarisation scans. This might be related to the existence of defects in the ZrO2 layer behaving as initiation sites for localised attack (pitting).

277 •

•

•

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The introduction of 2-mercaptobenzothiazole and cerium nitrate do not affect the morphology of sol–gel ZrO2 pre-treatments significantly. However, the possibility cannot be excluded that the addition of inhibitor to the sol–gel solution contributes to the formation of discontinuities in the pre-treatment. 2-Mercaptobenzothiazole produces a marked reduction in cathodic current density in potentiodynamic polarisation scans, in accordance with its well known ability to deactivate cathodic reactions. Moreover, the passive range exhibited by ZrO2 pre-treatment doped with 2-mercaptobenzothiazole confirms its good barrier properties. It shows a clear inhibition effect over a relatively short immersion time in electrolyte (about 24 h). The inhibition effect vanishes for longer immersion times probably due to complete leaching of inhibitor. This effect might be limited by deposition of a paint onto the ZrO2 pre-treatment. Nevertheless, the long-term corrosion resistance of ZrO2 pre-treatments doped with this inhibitor should be considered good taking into account the thin nature of the pre-treatment. Cerium nitrate does not clearly improve the electrochemical behaviour of ZrO2 pre-treatments. Further research is necessary before this inhibitor can be used in the coating system discussed in this work.

Acknowledgements This paper and the work it concerns were generated in the context of the MULTIPROTECT project, funded by the European Community as contract N° NMP3-CT2005-011783 under the 6th Framework Programme for Research and Technological Development. References 1. T. L. Metroke, R. L. Parkhill and E. T. Knobbe, Prog. Org. Coat., 41 (2001), 233. 2. R. Di Maggio, L. Fedrizzi and S. Rossi, J. Adhes. Sci. Technol., 15 (2001), 793. 3. M. Bethencourt, F. J. Botana, J. J. Calvino, M. Marcos and M. A. Rodriguez-Chacon, Corros. Sci., 11 (1998), 1803. 4. J. H. Osborne, Prog. Org. Coat., 41 (2001), 280. 5. M. Guglielmi, J. Sol-Gel Sci. Technol., 8 (1997), 443. 6. R. L. Ballard, J. P. Williams, J. M. Njus, B. R. Kiland and M. D. Soucek, Eur. Polym. J., 37 (2001), 381. 7. T. L. Metroke and A. Apblett, Prog. Org. Coat., 51 (2004), 36. 8. N. N. Voevodin, N. T. Grebasch, W. S. Soto, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 24. 9. A. M. Beccaria, G. Padeletti, G. Montesperelli and L. Chiaruttini, Surf. Coat. Technol., 111 (1999), 240. 10. A. M. Beccaria and L. Chiaruttini, Corros. Sci., 41 (1999), 885. 11. X. F. Yang, D. E. Tallman, V. J. Gelling, G. P. Bierwagen, L. S. Kasten and J. Berg, Surf. Coat. Technol., 140 (2001), 44. 12. N. N. Voevodin, V. N. Balbyshev and M. S. Donley, Prog. Org. Coat., 52 (2005), 28. 13. H. Yang and W. J. van Ooij, Prog. Org. Coat., 50 (2004), 140. 14. M. L. Zheludkevich, K. A. Yasakau, S. K. Poznyak and M. G. S. Ferreira, Corros. Sci., 47 (2005), 3368. 15. A. N. Khramov, N. N. Voevodin, V. N. Balbyshev and M. S. Donley, Thin Solid Films, 447–448 (2004), 549.

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16. M. L. Zheludkevich, R. Serra, M. F. Montemor, I. M. Miranda Salvado and M. G. S. Ferreira, Electrochim. Acta, 51 (2005), 208. 17. P. Campestrini, H. Terryn, A. Hovestad and J. H. W. de Wit, Surf. Coat. Technol., 176 (2004), 365. 18. N. N. Voevodin, N. T. Grebasch, W. S. Soto, F. E. Arnold and M. S. Donley, Surf. Coat. Technol., 140 (2001), 24. 19. V. Moutarlier, B. Neveu and M. P. Gigandet, Surf. Coat. Technol., 202 (2008), 2052. 20. J. H. Osborne, K. Y. Blohowiak, S. R. Taylor, C. Hunter, G. Bierwagon, B. Carlson, D. Bernard and M. S. Donley, Prog. Org. Coat., 41 (2001), 217. 21. L. Fedrizzi, R. Di Maggio, S. Rossi and L. Leonardelli, Benelux Metall., 43 (2003), 15. 22. F. Andreatta, P. Aldighieri, L. Paussa, R. Di Maggio, S. Rossi and L. Fedrizzi, Electrochim. Acta, 52 (2007), 7545. 23. P. Campestrini, E. P. M. van Westing, H. W. van Rooijen and J. H. W. de Wit, Corros. Sci., 42 (2000), 1853.

15 Influence of the doping agent on the corrosion protection properties of polypyrrole grown on Al-2024–T3 A. Balaskas, I. L. Danilidis, I. Kartsonakis and G. Kordas National Centre for Scientific Research, Demokritos, Agia Paraskevi, 15341, Attiki, Greece [email protected]

15.1

Introduction

Corrosion protection by conducting polymers has been under investigation for a long time [1–4], especially concerning the mechanism of protection. They are believed to act in many different ways, from facilitating the creation of a metal oxide film [5] to the release of a doping agent which can be properly chosen to act as an inhibitor [6]. Polyaniline has been of special interest [7–9] in contrast to polypyrrole [10] which is not widely used for corrosion protection applications. However, it is worth further investigation as it is chemically more stable and more easily deposited on a variety of substrates and at different pH values [11]. Electrochemical polymerisation of pyrrole gives rise to a homogeneous film carrying the electrolyte anions into the polymer backbone for reasons of electroneutrality [12,13]. Under corrosive conditions, the corrosion-generated current can pass through the polymer matrix forcing the release of the dopants. The proper choice of dopant provides special characteristics to the polymer. For corrosion protection purposes, these anions act as inhibitors. The extent of their inhibiting action depends on the kind of and concentration of the dopants and on the processing conditions. Cerium ions are very promising candidates for the replacement of toxic chromates. They provide dominant cathodic protection by forming an insoluble oxide on corrosion sites [14–17]. Cerium salts, especially the nitrates, have been used and studied extensively for corrosion protection. Furthermore, various organic molecules have been cited in the literature as potential corrosion inhibitors [18–21]. Oxalic acid (OA) forms an iron oxalate film on the conducting polymer film on iron [22] and carbon steel [23] substrates which passivates the surface. This has also been reported for polypyrrole films on aluminium [24] and aluminium alloy 2024 [25]. Most of the proposed systems are surfactants carrying a hydrophilic and a lipophilic part. Polyaniline doped with para-toluene sulphonic acid (pTSA) on iron substrates has been studied for its corrosion protection properties [26], and also combined with polypyrrole coatings to study their structural and thermal properties [27]. Phenyl phosphonic acid (PPA) has also been studied as a doping agent in polythiophene [28] and polypyrrole [29] providing anodic protection on mild steel; it has also been used to dope polyaniline coatings to protect iron substrates [26] and aluminium from sol–gel films [30]. Camphor sulphonic acid (CSA) has been used mainly in polyaniline systems and was found to increase conductivity 279

Influence of the doping agent on the corrosion protection properties

280

[31] and also to protect aluminium alloy 2024–T3 [32,33], iron [26] and other metals [34] from corrosion. This work focuses on the influence of the doping agent in polypyrrole on the corrosion protection of Al alloys. Cyclic voltammetry was used for film deposition, while open circuit potential and DC polarisation measurements took place at different time intervals to evaluate the corrosion protection properties of the film over time. SEM was employed to evaluate the quality of the coatings. 15.2 15.2.1

Experimental Chemicals

All chemicals were of analytical reagent grade and used as-received except pyrrole (ppy) which was distilled twice before use. Camphor sulphonic acids (CSA), p-toluene sulphonic acid (pTSA), phenyl phosphonic acid (PPA), oxalic acid (OA) and cerium nitrate (Ce(NO3)3) were all purchased from Sigma–Aldrich. All solutions were prepared with distilled deionised water before use. The molecular structures of the dopants are shown in Fig. 15.1. 15.2.2

Film growth

A three-electrode electrochemical cell was used for electropolymerisation connected to a Solartron 1470 MultiStat potentiostat. An Al-2024 panel, previously cleaned by etching in NaOH and desmuting in HNO3, was the working electrode and a platinum foil served as the auxiliary electrode. All potentials were recorded versus a saturated calomel electrode (SCE) serving as the reference electrode. A 0.2 M electrolytic solution was used in the case of CSA, pTSA, PPA and OA, with a 0.1 M solution of Ce(NO3)3 and 0.1 M pyrrole monomer for all solutions. Deposition took place under a potentiodynamic regime by sweeping the potential from –1 to +3 V vs SCE for five cycles at a scan rate of 30 mV s–1. 15.2.3

Open circuit potential measurement (OCP)

The potential at equilibrium was recorded during the first 6 h of immersion in 0.5 M NaCl to obtain an indication of the water uptake of the film and the passivation of the oxidisable metal surface covered with the conducting polymer [35]. A typical two-electrode cell was used for this purpose where the coated Al-alloy sample was the working electrode and a saturated calomel electrode (SCE) was used as the reference electrode. The cell was connected to a Solartron 1470 MultiStat potentiostat.

15.1 Molecular structures of the doping agents

281

Self-healing properties of new surface treatments

15.2.4

Corrosion tests by DC polarisation

DC polarisation tests took place in a three-electrode electrochemical cell where the area of the working electrode (modified Al-2024 T3) exposed to the corrosive solution was 1 cm2. For these measurements, a platinum foil and a SCE were used as the counter and the reference electrodes, respectively. Potentiodynamic tests were carried out by scanning the potential between –1 V and 0 V vs SCE at a rate of 1 mV s–1. The resulting graphics were fitted to the Tafel plots allowing the corrosion current and potential, and the polarisation resistance to be extrapolated. Corrware software was used for fitting. 15.3

Results and discussion

15.3.1 Polymer deposition by cyclic voltammetry and morphological characterisation Cyclic voltammetry was chosen for polymer deposition as a one-step process. Figure 15.2a presents a typical voltammogram obtained for polypyrrole electrodeposition in the presence of p-TSA as the dopant. In the first cycle, the strong passivation of Al is present and expressed by the peak at +1.5 V vs SCE. After the first cycle, the surface is passivated, partially by the polymer deposition and partially by the formation of the corresponding oxide which dramatically decreases the current intensity in the next cycles. Polypyrrole oxidation takes place at potentials close to 0 V vs SCE followed by a reduction at lower potentials. The peak appearing at –0.75 V vs SCE only within the first two cycles is attributed to an anion exchange reaction. Very similar cyclic voltammograms were obtained for coatings doped with each of the organic molecules (Fig. 15.2a–d). The same strong passivation peak and lower oxidation currents for the rest of the cycle were observed. In contrast, with nitrates as dopants, the passivation peak did not appear in the first cycle and there was no decrease in the anodic currents after the first cycle (Fig. 15.2e); this lack of passivation is probably due to the less acidic environment during the electropolymerisation process (pH 4.71) in contrast with the solutions containing organic acids, for polypyrrole p-TSA pH 1.00, polypyrrole PPA pH 1.38, polypyrrole CSA pH 1.01 and polypyrrole oxalic acid pH 1.16 at a temperature of 20.1oC. A phenomenon that possibly occurs for all coatings is the over-oxidation of polypyrrole resulting in the degradation of part of the coating as the formation potential of all films is driven to 3 V vs SCE [36,37]. The lower oxidation current after the first cycle for the organically doped coatings might be due to the over-oxidation of the polymer [38]; but the peak in the negative range of the potential is attributed to anion expulsion from the polypyrrole film, except in the case of PPA, showing that the main part of the polymer is in the conductive state. SEM images show the morphology of the resulting coating (Fig. 15.3). Typical polypyrrole morphology is obtained. This structure is reinforced by the presence of the p-TSA which is an amphiphilic surfactant. Using different doping anions, there are changes in surface morphological characteristics. When CSA is used to form the coating, the surface is rougher and with fewer defects than p-TSA, PPA and oxalic acid doped polymer coatings. The coatings doped with nitrates form a surface with similar roughness to p-TSA and PPA but with a reduced number of defects. There are also reports associating the hydrophilicity/hydrophobicity of the doping anions with the morphology of the films [39], suggesting that anions containing more hydrophilic

Influence of the doping agent on the corrosion protection properties

282

15.2 Electrodeposition diagrams (sweep rate 30 mV s–1: (a) pyrrole (0.1 M), p-TSA (0.2 M); (b) pyrrole (0.1 M), CSA (0.2 M); (c) pyrrole (0.1 M), PPA (0.2 M); (d) pyrrole (0.1 M), oxalic acid (0.2 M) and (e) pyrrole (0.1 M), cerium nitrate (0.1 M)

groups such as sulphates form rougher coatings than the less hydrophilic anions such as nitrates, in agreement with the present results. 15.3.2

Open circuit potential measurements, DC polarisation tests

The measurement of the open circuit potential (OCP) may provide information on the behaviour of the film while exposed to 0.5 M NaCl solution. Figure 15.4 presents a graph of OCP within the first 6 h of immersion to give an indication of the behaviour of the coating doped with different anions in the early period. From the electrodeposition diagrams and SEM images, the different properties of the film due to the different doping agents can be clearly seen. For all of the coatings, except that doped with OA, a decrease in the OCP is observed due to the water uptake of the film. For some films, the procedure lasts longer than for others, showing different properties for every film by changing the doping agent. When the dopants

283

Self-healing properties of new surface treatments

15.3 SEM images (a) polypyrrole p-TSA, (b) polypyrrole PPA, (c) polypyrrole CSA, (d) polypyrrole cerium nitrate, (e) pyrrole oxalic acid

are nitrates and physically absorbed cerium cations, the water uptake procedure is rapid and the potential stabilises at around –0.6 V vs SCE after 0.5 h of immersion; very similar OCP was observed with anions of p-TSA. The stronger influence of the electrolytic solution is observed for the coating doped with PPA where a stable state was reached after 1.5 h of immersion. In the case of OA and CSA, the OCP stabilises very quickly after only a few minutes of immersion; in particular, the OCP of the bare aluminium alloy and the coating formed with OA are very similar. The behaviour of the last two coatings is caused by the large number of defects in the surface of the coating as observed from SEM images.

Influence of the doping agent on the corrosion protection properties

284

15.4 Open circuit potential for the coatings in the early period of immersion

Figures 15.5 to 15.9 present the polarisation curves and the open circuit measurements obtained for each of the studied coatings at different times of immersion. It can be easily concluded that the behaviour of the polypyrrole film strongly depends on the doping agent. It must be underlined that the corrosion current density values cannot be regarded as the only criterion for evaluation of the coating. These films are mainly conducting which means that a considerable amount of the current density is related to the film conductivity and not only to the corrosion process. Figure 15.5a shows the film doped with nitrates using the corresponding cerium salt. According to theory, polypyrrole is mainly doped with nitrates [12] and cerium ions can be found physically absorbed into the polymer matrix. The response obtained after 48 h of immersion shows a lower corrosion current compared to that obtained at 24 h. This indicates a recovery with time and an improvement in the protection which can be due either to the formation of an underlying oxide or to the release of the dopants. The fact that the corrosion potential decreases in the same time interval suggests that the improvement is not consequential on oxide formation (barrier) alone as this would be accompanied by a potential increase due to stifling of metal dissolution, but rather due to dominant cathodic protection as the cathodic part of the curve decreases significantly. This is due to the release of nitrates and part of the contribution of the cerium cations. The passivation range increases in the same time interval, meaning that an anodic protection is also present. The open circuit potential measurement (Fig. 15.5b) shows a tendency for stabilisation of the potential after 24 h of immersion time in the corrosive environment. At 72 h of immersion, the coating shows a dominant barrier effect or anodic inhibition. The gradual shift of the corrosion potential to less electronegative values is due to the anodic current reduction, hence suppression of metal dissolution, that is observed from successive polarisation curves. The fitted values of the free corrosion potential, net corrosion

285

Self-healing properties of new surface treatments

15.5 (a) DC polarisation curves for ppy-Ce(NO3)3 at different immersion times; (b) ppy-Ce(NO3)3 open circuit potential measurement

Table 15.1 Corrosion current Icorr (A), free corrosion potential Ecorr (mV) and polarisation resistance for polypyrrole coating doped with cerium nitrate ppy-Ce(NO3)3 24h 48h 72h

Icorr (Acm–2)

Ecorr (mV/SCE)

Rp (kΩcm–2)

9.83E-06 1.84E-06 9.75E-07

–703.73 –748.56 –659.67

2.65 14.19 26.76

Influence of the doping agent on the corrosion protection properties

286

density and polarisation resistance, all relative properties of our system are shown in Table 15.1. Since the coating is going from the conducting to non-conducting state, the net current will be reduced and the inhibitor dopant will also cause the passivation and corrosion protection of the surface. Figure 15.6a shows the polarisation curves acquired after 24, 48 and 72 h of exposure to sodium chloride solution for the film grown in the presence of PPA. For this coating, the mechanism of protection is quite different. The film shows similar behaviour at 48 and 72 h with corrosion current and potential values very close introducing a rather stable and resistive coating. After an initial stable open circuit potential at

15.6 (a) DC polarisation curves for ppy-PPA at different immersion times; (b) ppy-PPA open circuit potential measurement

287

Self-healing properties of new surface treatments

approximately –600 mV (Fig. 15.6b) and fluctuation of the open circuit potential between 24 and 50 h of exposure, a relatively more stable behaviour at –780 mV is reached. Overall, there is a barrier effect but it starts to diminish after 24 h as can be concluded by comparing the results to those obtained after 48 h of immersion. The increase in the anodic current indicates that the film starts to degrade as can be seen from Fig. 15.6a and Table 15.2 where the fitted values of the free corrosion potential, net corrosion density and polarisation resistance are tabulated. The repeatability of the experiments is clear from Fig. 15.6b. The film doped with p-TSA preserves the protective properties for longer immersion times, reaching 72 h without significant degradation, as can be seen from Fig. 15.7a which shows polarisation curves acquired at 24, 48 and 72 h of exposure to sodium chloride solution. The corrosion potential (Fig. 15.7b) remains almost stable while the corrosion current decreases with time (Table 15.3). The cathodic protection is more evident here for this coating compared to the other pre-treatments. This stability through time may be explained by the nature of the dopant. Compared to the nitrates, p-TSA anions are of larger size making the release and complete reduction in the film slower, thus the protection longer. This also explains the decrease in the corrosion current from 24 to 48 h. A different behaviour was observed for the ppy-CSA coating (Fig. 15.8a,b). The open circuit potential is relatively unstable indicating the onset of metastable pitting corrosion from the early hours. The DC polarisation tests show that the corrosion potential is stable over time with a current increase after 72 h (Table 15.4). It is important to note that at 72 h, the corrosion current (both the anodic and cathodic current) has not significantly increased despite the fluctuation of the open circuit potential, whilst keeping the corrosion potential at a safe margin from the pitting potential. This suggests that the fluctuations of the open circuit potential are weak metastable pitting events. When oxalic acid is used to form the polypyrrole coating, the result is a film that is very stable in time without a significant change in the corrosion behaviour properties after 72 h of immersion to sodium chloride solution. The net corrosion current naturally increases after 72 h of immersion and is followed by a decrease in the polarisation resistance, but overall, it is a very stable system, with a safe passivation range (Fig. 15.9a,b, Table 15.5). Over-oxidation of a part of polypyrrole films would have occurred as discussed, but the amount of degraded polypyrrole should be small as the release of doping anions, a property of well formed conducting polymer film, is clearly seen. 15.4

Conclusion

Polypyrrole can be easily grown on AA 2024–T3 substrates in the presence of various electrolytes, from mineral salts to organic acids. This ease of deposition is a highly Table 15.2 Corrosion current Icorr (A), free corrosion potential Ecorr (mV) and polarisation resistance for polypyrrole coating doped with PPA ppy-PPA 24h 48h 72h

Icorr (Acm–2)

Ecorr (mV/SCE)

Rp (kΩcm–2)

5.20E-06 6.69E-06 7.26E-06

–639.08 –786.34 –762.93

5.02 3.90 3.59

Influence of the doping agent on the corrosion protection properties

288

15.7 (a) DC polarisation curves for ppy-pTSA coating at different immersion times; (b) ppy-pTSA open circuit potential measurement Table 15.3 Corrosion current Icorr (A), free corrosion potential Ecorr (mV) and polarisation resistance for polypyrrole coating doped with p-TSA ppy-pTSA 24h 48h 72h

Icorr (Acm–2)

Ecorr (mV/SCE)

Rp (kΩcm–2)

1.62E-05 4.65E-06 5.18E-06

–723.68 –745.93 –739.01

1.61 5.61 5.03

desirable property for the development of corrosion protective coatings. In this study, the chosen electrolytes were molecules known for their inhibiting properties. Their anticorrosion properties were evaluated by DC polarisation tests and open circuit potential measurements.

289

Self-healing properties of new surface treatments

15.8 (a) DC polarisation curve for ppy-CSA film at different immersion times; (b) ppy-CSA open circuit potential measurement

Table 15.4 Corrosion current Icorr (A), free corrosion potential Ecorr (mV) and polarisation resistance for polypyrrole coating doped with CSA ppy-CSA 24h 48h 72h

Icorr (Acm–2)

Ecorr (mV/SCE)

Rp (kΩcm–2)

7.14E-06 5.34E-06 1.19E-05

–707.50 –708.90 –708.16

3.65 4.89 2.20

The results showed that, depending on the doping agent, the polypyrrole-based coating can behave differently giving rise to different protection mechanisms. Thus, cerium salts give rise to well protecting coatings showing a progressively stronger barrier effect as well as cathodic current reduction, at longer exposure times in the

Influence of the doping agent on the corrosion protection properties

290

15.9 (a) DC polarisation curve for ppy-OA film at different immersion times; (b) ppy-OA open circuit potential measurement

Table 15.5 Corrosion current Icorr (A), free corrosion potential Ecorr (mV) and polarisation resistance for polypyrrole coating doped with oxalic acid ppy-OA 24h 48h 72h

Icorr (Acm–2)

Ecorr (mV/SCE)

Rp (kΩcm–2)

3.64E-06 4.84E-06 5.92E-06

–766.13 –824.29 –831.71

7.17 5.39 4.41

291

Self-healing properties of new surface treatments

corrosive medium, and possibly suggesting a self-healing effect. An increase in the cerium content could lead to longer protection times, but this is still being investigated. From the different organic acids used, the sulphonic derivates exhibit a rather similar behaviour in open circuit potential for the first 6 h which is probably related to the strong amphiphilic character of the molecule. When PPA is used, the coatings show a barrier effect for a short time and signs of possible self-healing with longer exposure; its anticorrosive properties are very similar to cerium nitrate. P-TSA, oxalic acid and CSA provide very stable and passive films in time, with slight deterioration of the last two at longer exposures. For all of the films tested, it was found that the release of the dopant does not occur instantly with immersion in the solution but rather progressively in time at different rates for each of the coatings. The recovery from corrosion activity by a tendency to re-stabilise the open circuit potential in combination with analysis of DC polarisation curves, suggests self healing. However, this needs to be investigated further. Consequently, polypyrrole is a promising material as it can provide a reservoir where different types of molecules can be attached to the backbone of the polymer. For corrosion protection applications, polypyrrole may be used as the inhibitor reservoir which can be released on demand. Acknowledgements We thank the European Integrated Project MULTIPROTECT (NMP3-CT-2005011783) for financial support. References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19. 20.

S. Patil, S. R. Sainkar and P. P. Patil, Appl. Surf. Sci., 225 (2004), 204. A. M. Fenelon and C. B. Breslin, Surf. Coat. Technol., 190 (2005), 264. K. Aramaki, Corros. Sci., 42 (2000), 1975. T. L. Nguyen, B. Garcia, C. Deslouis and L. Q. Xuan, Electrochim. Acta, 46 (2001), 4259. J. L. Camalet, J. C. Lacroix, S. Aieyach and P. C. Lacaze, J. Electroanal. Chem., 445 (1998), 117. S. Koehler, M. Ueda, I. Efimov and A. Bund, Electrochim. Acta, 52 (2007), 3040. S. B. Saidman and J. B. Bessone, J. Electroanal. Chem., 521 (2002), 87. R. Racicot, R. Brown and S. C. Yang, Synth. Met., 85 (1997), 1263. T. K. Rout, G. Jha, A. K. Singh, N. Bandyopadhyay and O. N. Mohanty, Surf. Coat. Technol., 167 (2003), 16. G. Paliwoda-Porebska, M. Stratmann, M. Rohwerder, K. Potje-Kamloth, Y. Lu, A. Z. Pich and H. J. Adler, Corros. Sci., 47 (2005), 3216. S. B. Saidman, J. Electroanal. Chem., 534 (2002), 39. Z. Gao, J. Bobacka, A. Lewenstam and A. Ivaska, Electrochim. Acta, 39 (1993), 755. P. A. Kilmartin and G. A. Wright, Electrochim. Acta, 43(21–22) (1998), 3091. A. Pepe, M. Aparcio, S. Cere and A. Duran, J. Non-Cryst. Solids, 348 (2004), 162. A. L. Rudd, C. B. Breslin and F. Mansfeld, Corros. Sci., 42 (2000), 275. D. R. Arnott, N. E. Ryan and B. R. W. Hinton, Appl. Surf. Sci., 22/23 (1985), 236. A. J. Aldykewicz, H. S. Isaacs and A. J. Davenport, J. Electrochem. Soc., 142(10) (1995), 3342. K. G. Conroy and C. B. Breslin, Electrochim. Acta, 48 (2003), 721. A. Kitani, K. Satoguchi, H. Q. Tang, S. Ito and K. Sasaki, Synth. Met., 69 (1995), 129. I. Mav, M. Zigon and A. Sebenik, Synth. Met., 101 (1999), 717.

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21. M. M. Osman, R. A. El-Ghazawy and A. M. Al-Sabagh, Mater. Chem. Phys., 80 (2003), 55. 22. C. B. Breslin, A. M. Fenelon and K. G. Conroy, Mater. Des., 26 (2005), 233. 23. R. Rajagopalan and J. O. Iroh, Electrochim. Acta, 46 (2001), 2443. 24. G. S. Akundi and J. O. Iroh, Polymer, 42 (2001), 9665. 25. M. A. Arenas, L. G. Bajos, J. J. de Demborenea and P. Ocon, Prog. Org. Coat., 62 (2008), 79. 26. G. Williams, A. Gabriel, A. Cook and H. N. McMurray, J. Electrochem. Soc., 153(10) (2006), B425. 27. A. Levent, J. Hacaloglu, L. Toppare and Y. Yagci, Synth. Met., 135 (2003), 457. 28. T. Tuken, B. Yazici and M. Erbil, Appl. Surf. Sci., 239 (2005), 398–409. 29. T. Tuken, B. Yazıcı and M. Erbil, Appl. Sur. Sci., 252 (2006), 2311. 30. M. Sheffer, A. Groysman, D. Starosvetsky, N. Savchenko and D. Mandler, Corros. Sci., 46 (2004), 2975. 31. P. Sonar, A. L. Sharma, A. Chandra, K. Muellen and A. Srivastava, Curr. Appl. Phys., 3 (2003), 247. 32. J. C. Seegmiller, J. E. Pereira da Silva, D. A. Buttry, S. I. Cordoba de Torresi and R. M. Torresi, J. Electrochem. Soc., 152(2) (2005), B45. 33. S. F. Cogan, M. D. Gilbert, G. L. Holleck, J. Ehlirch and M. H. Jillson, J. Electrochem. Soc., 147(6) (2000), 2143. 34. J. E. Pereira da Silva, S. I. Cordoba de Torresi and R. M. Torresi, Prog. Org. Coat., 58 (2007), 33. 35. N. K. Naoi, M. Takeda, H. Kanno, M. Sakakura and A. Shimada, Electrochim. Acta, 45 (2000), 3413. 36. I. Fernandez, M. Trueba, C. A. Nunez and J. Rieumont, Surf. Coat. Technol., 191 (2005), 134. 37. F. Beck, M. Dahlhaus and N. Zahedi, Electrochim. Acta, 37(7) (1992) 1265. 38. N. C. T. Martins, T. Moura e Silva, M. F. Montemor, J. C. S. Fernandes and M. G. S. Ferreira, Electrochim. Acta, 53 (2008), 4754. 39. S.-J. Choi and S.-M. Park, J. Electrochem. Soc., 149(2) (2002), E26.

16 Self-healing coating with fluoro-organic compound on zinc Akihiro Yabuki and Ryo Kaneda Graduate School of Engineering, Hiroshima University, 1-4-1, Kagamiyama, Higashi-Hiroshima, Hiroshima, 739-7852 Japan [email protected]

16.1

Introduction

Chromate conversion coatings have been widely applied as surface treatments for zinc-coated and galvanised steel used for applications such as automobile parts, building structures, home appliances, etc., since they have excellent anti-corrosion properties, are inexpensive, and the coating solution is easily prepared. Environmental concerns, however, have necessitated the reduction and discontinuation of this process in recent years. Trivalent chromate conversion coating has been applied as an alternative technology to hexavalent chromate conversion coating [1,2]. In addition, it has been reported that the addition of cerium, molybdic acid, phosphoric acid and colloidal silica to the solution for coating was effective as an alternative technology for chromate conversion coatings [3–10]. An important characteristic required for these types of coatings is that of self-healing, so that a film coated on the surface is automatically repaired if a defect occurs in the film. It is well known that the repair mechanism in chromate conversion coatings is due to the hexavalent chromium ion, which has high reactivity. There is a series of studies of the cerium ion by Hinton and Wilson, who reported that it is as effective as the chromium ion, as an inhibitor in solution [11]. The action of the cerium ion resembles that of the chromium ion, and CeO2 acts as a barrier film. When a defect is generated, the cerium ion in a film repairs it due to dissolution from the film and oxidation of the defect [12]. Sol–gel films have good adhesion to both metallic substrates and organic coatings. Furthermore, the incorporation of inorganic nanoparticles can be a way of inserting corrosion inhibitors and preparing inhibitor nanoreservoirs for self-healing pre-treatment [9,13–18]. The release of organic inhibitors from the hybrid sol–gel matrix can be described as a pH-dependent, triggered-release mechanism [17,18]. Some of the present authors conducted corrosion tests for an aluminium alloy coated with several types of polymer to prevent corrosion in a heat exchanger, and it was found that fluorine resin coating not only had excellent barrier properties, but self-healing properties as well [19]. In this study, a coating with a fluoro-organic compound on zinc was investigated to find an alternative technology to hexavalent chromate conversion coating. The barrier and self-healing properties of the coating were evaluated in sodium chloride solution. Several fluoro-organic compounds with different terminal groups and various numbers of carbon atoms were coated on a pure zinc sheet using the dip method. The barrier properties of the coatings were evaluated in sodium chloride 293

Self-healing coating with fluoro-organic compound on zinc

294

solution using the electrochemical impedance method. The optimum coating condition was determined by the concentration of the fluoro-organic compound in the solution and the pH of the solution. The change in the polarisation resistance of the coatings scratched by a rod with a diamond tip was monitored, and the surface appearance of the specimen was observed after the corrosion test to elucidate the self-healing properties of the coatings. 16.2 16.2.1

Experimental Coating preparation

Seven types of fluoro-organic compound, all with different terminal groups (COOH, OH, COF) and various numbers of carbon atoms (Table 16.1) were used to coat the zinc substrate. All but one of the fluoro-organic compounds were dissolved in deionised water at room temperature. A fluoro-organic compound of F-OH(C5) was dissolved in deionised water at 50°C, because it could not be dissolved at room temperature. Concentrations of the fluoro-organic compounds in the solutions ranged from 100 to 5000 ppm, and the pH of the solutions was controlled in a range of 2–10 using phosphoric acid and sodium hydroxide. A zinc sheet of 99.99% purity was used as the substrate for coating. The dimensions of the specimens, cut from the zinc sheet, were 12 mm × 12 mm × 1 mm. The specimens were polished using a series of #2000 emery papers, followed by thorough rinsing in acetone with an ultrasonic bath and air drying. The polished specimen was immersed in the prepared solution with fluoro-organic compounds for 120 s, then it was rinsed in deionised water and air dried. A bare zinc specimen and a specimen treated in a 20 000 ppm chromate solution were used for reference. 16.2.2

Evaluation of barrier properties for coating

The barrier properties of the specimen coated with fluoro-organic compounds were evaluated in corrosive solution using electrochemical impedance and polarisation measurements. A 0.5% NaCl solution, maintained at 40°C, was used as the corrosion test liquid. The solution was air-saturated with an air pump. The specimen was put in a holder made of methylmethacrylate and connected to a conductive wire for electrochemical measurement. The measurement area of the specimen had a diameter of 5 mm. The electrochemical impedance and polarisation measurements were taken after the corrosion potential of the specimen became constant in the NaCl Table 16.1 Fluoro-organic compounds used for coating Symbol

Fluoro-organic compound

F-COOH(C5) F-COOH(C8) F-OH(C5) F-OH(C7) F-OH(C10) F-COF(C12) F-COF(C15)

(CF3)2C(CH3)COOH F(CF2)7COOH (CF3)2C(CH3)CH2OH H(CF2)6CH2OH F(CF2)8CH2CH2OH F(CF2)3O[CF(CF3)CF2O]2CF(CF3)COF F(CF2)3O[CF(CF3)CF2O]3CF(CF3)COF

Molecular weight 210.06 414.05 196.09 332.08 464.11 664.06 830.08

295

Self-healing properties of new surface treatments

solution. The exposure time was more than 10 min. The alternating current was superimposed on the corrosion potential before starting the electrochemical impedance measurement. The impedance of the specimen in the corrosive solution was measured using platinum counter and Ag/AgCl reference electrodes connected to a potentiostat (HA-150, Hokuto Denko Co.), a frequency response analyser (5010A, NF Co.), and a personal computer. Sine wave voltages (10 mV r.m.s.) at frequencies from 20 kHz through 10 mHz were superimposed on a given electrode potential. A computer software program was used to control the measurements through a General Purpose Interface Bus (GPIB). The difference in impedances measured at low and high frequencies was used to determine the polarisation resistance, since the phase shift was almost zero in the low and high frequency ranges. The polarisation curve of the specimen was measured in a 0.5% NaCl solution using the same system as for electrochemical impedance measurement. After the corrosion potential of the specimen immersed in the test solution had become almost constant, cathodic and anodic polarisation curves were measured at a sweep rate of 20 mV/min. Measurement data were inputted to a computer through a GPIB interface. The infrared reflection/absorption spectrum of the coatings was obtained using a Perkin-Elmer FT-IR Spectrum 100 spectrometer. The coating spectrum was ratioed against the spectrum of the substrate background. A pure fluoro-organic compound spectrum was obtained on insertion in the KBr pellets in the transmittance mode. A Joel JSM-6340F scanning electron microscope was used for observation of the coatings. 16.2.3

Evaluation of self-healing properties of the coating

A scratch was created on the coated specimen using a rod with a diamond tip on a scratch tester (IMC-1552, Imoto Machinery Co. Ltd.). The load for the scratch was 30 g. The scratched specimen was put in a holder and connected to a conductive wire, followed by immersion in a 0.0005 M NaCl corrosive solution, maintained at 40°C. The electrochemical impedance spectroscopy of the scratched specimen was measured at intervals of 4 h, or more, for a total of 24 h. Measurement data were analysed to calculate the polarisation resistance after the test. Scratched specimens were also immersed in a 0.0005 M NaCl corrosive solution at 40°C for 7 days. The surface appearance of the specimens was observed after the corrosion test to confirm the selfhealing properties of the coatings. The scratched part of the specimen was analysed after the corrosion test by Electron Probe Micro-Analysis (EPMA). 16.3 16.3.1

Results and discussion Barrier properties of coating

Electrochemical impedance spectroscopy of a bare zinc specimen and of specimens coated with fluoro-organic compounds F-COOH(C5) and F-COF(C12) at a concentration of 1000 ppm in solution are shown as a Nyquist plot in Fig. 16.1. The plots for the specimens were almost semicircular or slightly semi-oval. The difference in impedance value at high and low frequency, that is, the diameter of this depressed semicircle, corresponds to the resistance of the coating to corrosion. All specimens

Self-healing coating with fluoro-organic compound on zinc

296

16.1 Cole–Cole plot of specimens coated with two types of fluoro-organic compound with 1000 ppm concentration in solution, and a bare zinc specimen

coated with fluoro-organic compounds showed similar plots. The difference in impedance value at high and low frequencies was applied as an index for the barrier properties of the coating. The polarisation resistance of the specimen coated with various fluoro-organic compounds is shown in Fig. 16.2. The solution used for the coating contained 1000 ppm of the fluoro-organic compound. Comparison plots for chromate conversion coating and for the bare zinc specimen are also shown in Fig. 16.2. The polarisation resistances of the specimen coated with fluoro-organic compounds F-COOH(C5), F-COOH(C8) and F-OH(C7) were smaller than that of pure zinc, which resulted in inferior barrier properties. The polarisation resistances of specimens coated with fluoro-organic compounds F-OH(C5), F-OH(C10) and F-COF(C12) were slightly higher than that of bare zinc. These coatings had barrier properties, but they were quite low compared to chromate conversion coating. The polarisation resistance of the specimen coated with the fluoro-organic compound F-COF(C15) was quite high – equal to that of the chromate conversion coating. The barrier resistance of coatings increased with increasing number of carbon atoms in the compound, except for fluoro-organic compound F-OH(C5). The barrier properties of F-OH(C5) might be due to its molecular structure, which was non-linear, as described in Table 16.1. It was difficult to adsorb the fluorine atom in the fluoro-organic compound onto the zinc substrate, because the bond between the carbon atom and the fluorine atom is very strong. The carbon atoms of the fluoro-organic compounds used in this study were almost entirely covered by the fluorine atoms, so it was thought that the COOH, OH and COF terminals of the compounds interacted with the zinc surface. The fluoro-organic compounds F-COF(C12) and F-COF(C15), which also have nonlinear structures, might be adsorbed, thereby increasing density, which indicates a coating with excellent barrier properties.

297

Self-healing properties of new surface treatments

16.2 Polarisation resistance of specimens coated in solutions containing various fluoro-organic compounds with 1000 ppm concentration, chromate conversion coating, and a bare zinc specimen

16.3.2

Optimum conditions for coating with a fluoro-organic compound

Coating with the fluoro-organic compound F-COF(C15) was carried out under various conditions to determine the optimum conditions, defined by the concentration of the compound in the solution and the pH of the solution. Figure 16.3 shows the relationship between the polarisation resistance of the coating and the concentration of the fluoro-organic compound in a pH 3.2 solution. The filled circles show the mean values. The dotted lines in the figure are the polarisation resistances of the chromate conversion coating and the zinc substrate. The polarisation resistance of the specimen coated at a concentration of 500 ppm was almost the same as that of the zinc substrate. The polarisation resistance of the coating increased with increased concentration of the fluoro-organic compound in solution, and it showed a maximum value of 14 000 Ω cm2 at a concentration of 2000 ppm, which was larger than that of the chromate conversion coating. There was a wide variation in polarisation resistance at a concentration of 1000 ppm; however, the variation in polarisation resistance became smaller at concentrations higher than 2000 ppm. The amount of fluoroorganic compound that was sufficient to cover the surface of the specimen was 1000 ppm, but with this amount, some defects might occur in the coating due to coating conditions, since the resistance variation attributed to coating reproduction was observed. At concentrations higher than 2000 ppm, the amount of fluoro-organic compound was sufficient to cover the entire surface, even when the coating conditions, such as temperature and humidity in the drying process, were varied. A slight decrease in the polarisation resistance at high concentrations might be due to the formation of a reactive site by the fluoro-organic compound. Figure 16.4 shows the relationship between the polarisation resistance of the coated specimen and the pH of the fluoro-organic compound solution at a concentration of 2000 ppm. The polarisation resistances of the chromate conversion coating and the zinc substrate are shown in the figure as dotted lines. The polarisation resistance of

Self-healing coating with fluoro-organic compound on zinc

298

16.3 Effect of the concentration of fluoro-organic compound F-COF(C15) in pH 3.2 solution on the polarisation resistance of the coating

16.4 Effect of the pH of solution with fluoro-organic compound F-COF(C15) with 2000 ppm concentration on the polarisation resistance of the coating

the specimen coated in the solution at pH 2 and 10 was as low as it was with the zinc substrate. However, it was quite high at pH 5, compared to the chromate conversion coating. Observation of the solution showed it to be cloudy at pH 2 to 5, but it was transparent at pH 10. This indicated that the molecular state of the compound in the solution was different at each pH, that is, the fluoro-organic compound was agglomerated at low pH, and it became isolated in the solution with increasing pH. It seems

299

Self-healing properties of new surface treatments

that a uniform film did not form because the agglomeration of the compound was too high for coating at pH 2. The agglomeration of the compound in the solution at pH 5 was suitable for the formation of a barrier film. A barrier film no longer formed at pH 10, since the resistance was almost equal to that of the zinc substrate. As a result, an excellent barrier coating with the fluoro-organic compound F-COF(C15) could be formed in the solution containing the compound at a concentration of 2000 ppm at pH 5. The polarisation resistance of the coating greatly exceeded that of the chromate conversion coating. Using mass measurement of the specimen before and after the coating process, an attempt was made to determine the amount of fluoro-organic compound F-COF(C15) coated on the zinc substrate. However, since mass loss occurred, this could not be determined. This indicated that the zinc substrate was slightly corroded and the thickness of the coating was very small. Figure 16.5 shows the morphology of the specimen coated with the compound under these conditions. A thin film along the polished line was observed on the zinc substrate, and particles of 20 to 50 nm in diameter were involved in the film, resulting in a coating thickness of between 20 and 50 nm. The infrared reflectance spectrum of zinc coated with fluoro-organic compound F-COF(C15) at 2000 ppm is shown in Fig. 16.6a. For comparison, the same figure shows the infrared reflectance spectrum of the fluoro-organic compound F-COF(C15) in KBr pellets (Fig. 16.6b). Both spectra clearly correspond to the same compound, frequencies of 1143 to 1310 cm–1 are assigned to –C–F– and –CF2– stretching vibrations. Frequencies of 1776 and 1886 cm–1 in the fluoro-organic compound F-COF (C15) are assigned to the C=O stretching vibration [20]. But when the fluoro-organic compound F-COF(C15) is coated onto the zinc substrate (Fig. 16.6a), the frequencies shift lower and broaden. This implies the adsorption of the COF terminal of the compound on the zinc surface, although the IR technique used was neither sufficiently selective nor sensitive enough to provide information about the link between the zinc and the fluoro-organic compound at the coating–zinc alloy interface.

16.5 SEM micrograph of zinc coated with fluoro-organic compound F-COF(C15) with 2000 ppm concentration

Self-healing coating with fluoro-organic compound on zinc

300

16.6 IR spectrum of (a) zinc coated with fluoro-organic compound F-COF(C15) with 2000 ppm concentration; (b) fluoro-organic compound F-COF(C15)

Measurement of polarisation curves was carried out to examine the barrier effect of the fluoro-organic compound F-COF(C15). The polarisation curves for the specimen coated with the fluoro-organic compound F-COF(C15) at a concentration of 2000 ppm, and for bare zinc substrate, are shown in Fig. 16.7. The cathodic polarisation of the coated specimen was almost the same as that of the bare zinc substrate. On the other hand, the anodic polarisation curve of the coated specimen was different from that of the zinc substrate, that is, a passivation region was observed up to –0.4 V. The anodic current for the F-COF(C15) coating in the region was lower than that of bare zinc, so that the anodic reaction was greatly inhibited by the film formed on the zinc surface. The favourable barrier effect of the coating with the fluoro-organic compound was caused by inhibition of the anodic reaction. 16.3.3

Self-healing properties of coating

To examine the self-healing properties of the coating, a scratch was applied to the coating of the fluoro-organic compound F-COF(C15), which has excellent barrier properties, and then polarisation resistance was measured using the electrochemical impedance method monitoring a 0.0005 M NaCl solution for 24 h. The polarisation resistance of the scratched specimen with the applied chromate conversion coating and of the bare zinc substrate were also monitored in the corrosion solution. The time dependence of the polarisation resistance of the scratched specimen is shown in Fig. 16.8. The resistance of the zinc substrate decreased after immersion and reached half of the initial resistance after a 12 h immersion, then became constant. This indicates that anodic dissolution proceeded in the solution. The polarisation resistances of the scratched specimen coated with the fluoro-organic compound F-COF(C15) and the chromate conversion coating both decreased in the early stage of immersion.

301

Self-healing properties of new surface treatments

16.7 Polarisation curves for specimen coated with fluoro-organic compound F-COF(C15) with 2000 ppm concentration and bare zinc substrate

16.8 Change in polarisation resistance of scratched specimens

Self-healing coating with fluoro-organic compound on zinc

302

However, these values began to increase after a 10 h immersion, resulting in a resistance higher than the initial resistance after 24 h. Based on the well-known self-healing properties of the chromate conversion, an increase in the polarisation resistance implies the existence of self-healing properties. Therefore, it was confirmed that the coating with the fluoro-organic compound also has self-healing properties. In addition, a 7-day corrosion test was conducted with the scratched specimen to determine the self-healing effect of these coatings. The surface appearance of the scratched specimen after the corrosion test is shown in Fig. 16.9. The scratched part of the specimen coated with the fluoro-organic compound F-COF(C15) was not an initiation site for corrosion, although corrosion pits were found on the surface of the specimen. The scratched specimen of the chromate conversion coating was corroded over almost the entire surface after the 7-day corrosion test, and white rust was observed on the surface. White rust was also observed on the surface of the zinc specimen. The scratch applied to the zinc specimen was no longer visible. The rust was thicker than that formed on the chromate conversion coating. A backscattered electron (BSE) image of the specimen surface and line analysis of Zn and F between A and B across the scratch on the zinc surface coated with the

16.9 Surface appearance of scratched specimens tested in 0.0005 M NaCl solution for 7 days

303

Self-healing properties of new surface treatments

16.10 A BSE photograph of the specimen surface and line analysis of Zn and F between A and B across the scratch for the zinc surface coated with fluoro-organic compound F-COF(C15), after the surface was scratched and immersed in 0.0005 M NaCl for 7 days

fluoro-organic compound F-COF(C15), after immersion in 0.0005 M NaCl for 7 days, is shown in Fig. 16.10. The X-ray intensity of F seemed to increase slightly at the scratch, compared with that of the coating, indicating migration of the fluoroorganic compound F-COF(C15) by diffusion from the scratched film into the scratched surface. The X-ray intensity of Zn at the scratch was almost the same as that of the coating, because the scratch surface was covered with the fluoro-organic compound. If it is assumed that the fluoro-carbon compound diffused to the scratched part, the coating must not be polymeric, but could instead consist of agglomeration of the fluoro-organic compound, which could be easily desorbed and diffused by a change of environment. Since the thickness of the coating is assumed to be 20 to 50 nm, as shown in Fig. 16.6, and the length of the molecules is approximately 2 nm, there could be 10 to 25 layers of the molecule in the coating. A fluoro-carbon compound nearest the zinc substrate was adsorbed with the terminal C=O of the compound, and the compound in the coating possibly diffused and adsorbed onto the scratched zinc substrate. Increasing the pH near the zinc substrate brought about the release of the fluoro-organic compound [17,18].

Self-healing coating with fluoro-organic compound on zinc

304

16.11 Schematic representation of self-healing in a defect from the release of fluoro-organic compound by cathodic reaction due to increasing pH near the defect

A schematic representation of the self-healing effect in the fluoro-organic compound coating is shown in Fig. 16.11. The self-healing effect appears to be due to the inhibiting action of the fluoro-organic compound on the corrosion process. As a defect occurred in the coating, bare zinc was exposed to the corrosive solution, and then anodic dissolution occurred. The anodic process led to the generation of metal cations. Zn → Zn2+ + 2e– Negatively charged OH– ions were generated near the defect due to the following cathodic reaction. O2 + 2H2O + 4e– → 4OH– Moreover, a local increase in pH near the defect can promote the release of the fluoro-organic compound from the coating near the defect, because the compound was easily isolated at a high pH, as described in the effect of pH on the polarisation resistance shown in Fig. 16.4. The released fluoro-organic compound diffused gradually on bare zinc to form a barrier film, which resulted in the repair of defects. 16.4

Conclusions

Corrosion tests on specimens were conducted with various fluoro-organic compounds to prevent the corrosion of zinc, and the following results were obtained. 1) The barrier properties of the fluoro-organic compound F-COF(C15) were almost equal to those of the chromate coating. The optimum conditions for formation of a superior barrier coating of the fluoro-organic compound were a concentration of 2000 ppm and pH 5. 2) The coating with the fluoro-organic compound F-COF(C15) on pure zinc has self-healing properties, demonstrated by electrochemical measurements on the scratched specimen in the corrosion solution and observation of the surface appearance of the scratched specimen after the corrosion test. The fluoroorganic compound has self-healing properties superior to those of the chromate conversion coating.

305

Self-healing properties of new surface treatments

3) The self-healing effect of the coating appears to be due to the release of the fluoroorganic compound by the pH increase caused by the cathodic reaction in the corrosion process and the formation of a film on the defect. Acknowledgement This research was partially supported by a grant from the ‘JFE 21st Century Foundation’. References 1. H. Noguchi, J. Yoshino and Y. Matsuda, J. Surf. Finish. Soc. Jpn., 51(8) (2000), 856. 2. T. Aoe, J. Surf. Finish. Soc. Jpn., 49(3) (1998), 221. 3. S. Dalbin, G. Maurin, R. P. Nogueira, J. Persello and N. Pommier, Surf. Coat. Technol., 194 (2005), 363–371. 4. A. A. O. Magalhaes, I. C. P. Margarit and O. R. Mattos, J. Electroanal. Chem., 572 (2004), 433–440. 5. R. G. Buchheit, H. Guan, S. Mahajanam and F. Wong, Prog. Org. Coat., 47 (2003), 174–182. 6. Y. K. Song and F. Mansfeld, Corros. Sci., 48 (2006), 154–164. 7. V. Palanivel, Y. Huang and W. J. V. Ooij, Prog. Org. Coat., 53 (2005), 153–168. 8. M. L. Zheludkevich, R. Serra, M. F. Montemor and M. G. S. Ferreira, Electrochem. Commun., 7(8) (2005), 836–840. 9. M. L. Zheludkevich, R. Serra, M. F. Montemor, K. A. Yasakau, I. M. M. Salvado and M. G. S. Ferreira, Electrochim. Acta, 51(2) (2005), 208–217. 10. W. Trabelsi, P. Cecilio, M. G. S. Ferreira and M. F. Montemor, Prog. Org. Coat., 54 (2005), 276–284. 11. B. R. W. Hinton and L. Wilson, Corros. Sci., 29(8) (1989), 967–975. 12. R. G. Buchheit, S. B. Mamidipally, P. Schmutz and H. Guan, Corrosion, 58(1) (2002), 3–14. 13. M. L. Zheludkevich, R. Serra, M. F. Montemor, K. A. Yasakau, I. M. M. Salvado and M. G. S. Ferreira, Electrochim. Acta, 51(2) (2005), 208–217. 14. V. Sviatlana, V. Lamakaa, M. L. Zheludkevich, K. A. Yasakaua, M. F. Montemor, P. Cecílio and M. G. S. Ferreira, Electrochem. Commun., 8(3) (2006), 421–428. 15. M. L. Zheludkevich, L. M. Salvado and M. G. S. Ferreira, J. Mater. Chem., 15(48) (2005), 5099–5111. 16. S. V. Lamaka, M. L. Zheludkevich, K. A. Yasakau, R. Serra, S. K. Poznyak and M. G. S. Ferreira, Prog. Org. Coat., 58(2–3) (2007), 127–135. 17. D. G. Shchukin, M. L. Zheludkevich, K. Yasakau, S. Lamaka, M. G. S. Ferreira and H. Mohwald, Adv. Mater., 18 (2006), 1672–1678. 18. A. J. Vreugdenhil and M. E. Woods, Prog. Org. Coat., 53(2) (2005), 119–125. 19. A. Yabuki, H. Yamagami and K. Noishiki, Mater. Corros., 58(7) (2007), 497–501 20. D. Saikia and A. Kumar, Eur. Polym. J., 41(3) (2005), 563–568.

INDEX

Index Terms

Links

A AA 2024-T3 alloy

184

Abbott-Firestone curve

224

acetylacetone

199

active corrosion protection optimisation of active pigments protective properties of

17

220

230

234 1

9

3

alkaline etching

264

aluminium alloys, corrosion of

108

119

134

145

150

161

167

172

184

204

207

217

220

262

279

293

287

anion-exchange pigments

26

anionic inhibitors

173

anodic films

134

aqueous electropolymer-isation

239

atomic force microscopy

255

145

B barrier effects

5

41

172

291

293

300

benzotriazole

10

bio-inhibition

157

This page has been reformatted by Knovel to provide easier navigation.

Index Terms

Links

C camphor sulphonic acid (CSA)

279

carboxylic acids

173

catechol, use of

240

cathodic polarisation cathodic reactions

291

243

256

154

170

177

221

228

268

273

277

287

304

6

cation-exchange solids

24

cerium-based coatings

46

59

167

170

177

184

195

204

216

221

230

232

235

262

273

277

279

291

293

1

11

40

79

108

119

134

149

152

166

184

203

205

220

232

262

characterisation techniques chromates, use of and alter-natives to

223

293 conducting polymer coatings

17

copolymerization

246

crevice corrosion

163

259

critical pigment volume con-centration (CPVC) cyclic voltammetry cyclodextrins

256 281 21

D delamination of coating materials

198

This page has been reformatted by Knovel to provide easier navigation.

Index Terms

Links

E eco-friendly inhibitors, development of

151

efficiency of protection against corrosion

160

electroactive conducting polymers

238

electrochemical impedance spectroscopy (EIS)

67

136

199

213

221

225

230

242

255

260

274

295

electrochemical impedance spectroscopy (EIS) electrochemically determin-ed corrosion rate electronically conducting polymers (ECPs)

158 79

elemental mapping

226

encapsulation

155

F filiform corrosion

163

fluoro-organic compounds

293

optimum conditions for coating with

167

297

G glow discharge optical emiss-ion spectroscopy (GDOES)

264

271

276

H halloysite nanotubes

34

hectorite

16

hybrid films and coatings

18

228

105

203

hybrid polymers

208

211

217 hydrogen peroxide hydrotalcite

171 26

This page has been reformatted by Knovel to provide easier navigation.

Index Terms

Links

I impedance modulus, meas-urement of

274

inductively coupled plasma atomic emission spectro-scopy (ICPAES) inhibitor activity parameter inhibitor development

163 154 4

eco-friendly

151

factors required for success in

166

inhibitor selection, screening techniques for

157

inteferometry

224

intermetallic particles

175

ion-exchange structures

167

149

262

228

L lanthanides layer-by-layer (LbL) assem-bled shells leaching of inhibitors

204 30

43

163

198

277

236

271

M mercaptobenzothiazole (MBT) micaceous iron oxide microemulsion polymer-isation

27 6 27

microfluidics

161

morphological studies

226

multi-electrode arrays

158

multifunctional corrosion inhibitors

157

multilayer coatings

212

281

172

This page has been reformatted by Knovel to provide easier navigation.

Index Terms

Links

N nanocomposite layers

112

116

nanocontainers

21

30

nanoparticles

52

74

145

204

220

227

234

262

Nissan

14

O open circuit potential (OCP) measurement

280

287

organic coatings

148

220

anticorrosive pigments in incorporation of inhibitors into

1 165

organosilane-based treat-ments

41

79

173

osmotic blistering

17

168

204

oxalic acid

279

283

oxidation reactions

171

oxide nanoparticles

23

oxygen deprivation

154

P Paint Research Association paint systems

passivation

152 14

65

173

262

3

154

passive protection against corrosion

220

phenyl phosphonic acid

279

phosphatic coatings

167

pigment volume concent-rations (PVCs)

239

pitting scans plasma deposition process

149

153

286

260

93 222

This page has been reformatted by Knovel to provide easier navigation.

Index Terms

Links

plasma polymerisation (PP)

27

221

228

236

polarisation resistance and polarisation curves

72

172

224

269

272

284

295

135

239

279

173

238

279

153

169

177

7

149

poly(phenylenesulphide) (PPS)

15

polyacetylene

238

polyaniline (PANI)

120

polymer composites

12

polymerization

245

see also aqueous electro polymerisation; co-polymerisation; micro emulsion polymerisation polypyrrole (PPy)

135

potential transient measure-ments

96

potentiodynamic polaris-ation

72

Pourbaix diagram

171

protuberances

235

pyrrole-based silane (SiPy)

272

79

R rare earths, use of

150 184

red lead oxide reference electrodes reflow effect

1 158 12

release kinetics

168

resistance inhibition

154

Rutherford back-scattering spectroscopy (RBS)

176

This page has been reformatted by Knovel to provide easier navigation.

Index Terms

Links

S salt spray testing (SST)

131

136

145

187

221

228

235

242

258

264

276

281

198 scanning electron micro-scopy (SEM)

299 scanning vibrating electrode technique (SVET)

28

72

3

7

149

173

199

207

217

220

225

230

232

235

262

275

289

293

295

300

definitions of

11

42

self-sealing coatings

11

207

217

39

59

65

77

221

262

self-healing coatings

based on active protection

17

based on nano-/microcon-tainers

20

silane coatings

anti-corrosion properties of

41

formation of

40

with active ions

43

with active species

42

silica coatings

202

217

siloxane coatings

220

226

230

232

39

43

105

119

185

199

202

209

217

262

276

293

sol-gel technologies

strontium chromate superprimers surface analysis techniques

149 19 173

This page has been reformatted by Knovel to provide easier navigation.

Index Terms

Links

T thiazole compounds

262

topographical grain analyses

224

triazole compounds

262

triggered release of corrosion inhibitors

151

V vibrational techniques

173

176

see also scanning vibrating electrode technique

X X-ray photoelectron spectro-scopy (XPS)

248

Z zeolite particles

24

zinc, coatings on

293

zinc chromate

149

zinc phosphate

165

172

zinc-rich primers

1

7

Zr-based coatings

46

262

ZrO2 layers

262

This page has been reformatted by Knovel to provide easier navigation.