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

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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.


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


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


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


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


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


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


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]


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


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


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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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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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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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


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


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


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]

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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


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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


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]


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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].

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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)


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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


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


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


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


34

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]

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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


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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

The use of modified nanostructured and functional organosilane coatings 44

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


46

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


48

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

49


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


50

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.


52

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


54

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


56

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 (∅ 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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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

81


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.


82

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

83


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


84

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

85


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


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

87


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)


88

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

89


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


90

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

91


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


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 (---)

92

93


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


94

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

95


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


96

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

97


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


98

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

99


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


100

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)

101


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.


102

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

Sol–gel derived hybrid materials as functional coatings

106

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


110

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 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

169


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

Advances in the selection and use of rare-earth-based inhibitors

170

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].

171


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.


172

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


174

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.

175


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


176

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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181 89. 90. 91. 92. 93. 94. 95. 96. 97. 98. 99. 100.

101. 102. 103. 104. 105. 106. 107. 108. 109.

110.

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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

185


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

Corrosion inhibiting cerium compounds

186

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

187


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


188

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

189


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


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

191


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,


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

193


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%

195


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


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

197


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.


198

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

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