Adhesion aspects of polymeric coatings. Vol. 2

Adhesion Aspects of Polymeric Coatings, Volume 2 K.L. Mittal, Editor VSP Adhesion Aspects of Polymeric Coatings, Vol...

Author: K. L. Mittal

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Adhesion Aspects of Polymeric Coatings, Volume 2

K.L. Mittal, Editor

VSP

Adhesion Aspects of Polymeric Coatings, Volume 2

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ADHESION ASPECTS OF POLYMERIC COATINGS VOLUME 2

Editor: K.L. Mittal

UTRECHT BOSTON, 2003

VSP BV P.O. Box 346 3700 AH Zeist The Netherlands

Tel: +31 30 692 5790 Fax: +31 30 693 2081 [email protected] www.vsppub.com

© VSP BV 2003 First published in 2003 ISBN 90-6764-377-7

All rights reserved. No part of this publication may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, recording or otherwise, without the prior permission of the copyright owner.

Printed in The Netherlands by Ridderprint bv, Ridderkerk

Contents

Preface Interphase: Formation, characterization and relevance to practical adhesion A.A. Roche, J. Bouchet and S. Bentadjine

vii

1

Depletion, a key factor in polymer adhesion G. Frens

21

Attaining adhesion/cohesion within painted plastics R.A. Ryntz

29

Scanning electric potential microscopy (SEPM) and electric force microscopy (EFM) imaging of polymer surfaces E.F. de Souza, M.M. Rippel, A.J. Keslarek, A. Galembeck, C.A.R. Costa, É.T. Neto and F. Galembeck

45

The residue (smut) formed on aluminum alloys during hydrofluoric acid etching and its effect on a coating process A.P.S. Tihaiya, J.P. Bell and G.D. Davis

65

Surface modification of metals by silanes D.Zhu and W.J. van Ooij

81

Application of X-ray photoelectron spectroscopy in assessing the adsorption of siloxane polymers onto E-glass fibers L.G. Britcher, D. Kehoe and J.G. Matisons

101

Surface modification of polyphenylene sulfide plastics to improve their adhesion to a dielectric adhesive Y. Wang and S. Rak

121

Metal surface conditioning concepts for resin bonding in dentistry P. Pfeiffer, I. Nergiz and M. Özcan

137

vi

Contents

Measurement of internal stresses in polymeric coatings using time resolved fluorescence T. Ikawa and T. Shiga

151

Adhesion of an alkyd paint to cold rolled steel sheets: Effect of steel surface composition S. Maeda

165

Analysis of the wet adhesion of coatings on wood M. de Meijer

183

Modified tape test: Measurement of adhesion of insulator films to low dielectric constant organic polymers L.L.N. Goh, S.L. Toh, S.Y.M. Chooi, Y. Xu and T.E. Tay

203

Adhesion Aspects of Polymeric Coatings, Vol. 2, pp. vii–viii Ed. K.L. Mittal © VSP 2003

Preface This volume documents the proceedings of the Second International Symposium on Adhesion Aspects of Polymeric Coatings held under the auspices of MST Conferences in Newark, New Jersey, May 25–26, 2000. The premier symposium on this topic was held under the aegis of the Electrochemical Society in Minneapolis, Minnesota in 1981, the proceedings of which were properly documented in a hard-bound book [1]. As almost 20 years had passed since the first symposium was held so we decided to organize the second event on this topic. In the interim, there had been a great deal of research activity relative to the adhesion aspects of polymeric coatings, so this symposium was both timely (rather overdue) and needed. Polymeric coatings are used for a variety of purposes, e.g., decorative, protective, functional (as dielectrics or insulators) and a special application of polymeric coatings is their use as lithographic materials for making integrated circuit elements. Irrespective of the intended purpose of the coating, it must adequately adhere to the underlying substrate, otherwise delamination and other undesirable phenomena can occur. So the need to understand the factors which influence adhesion of polymeric coatings and to control it to a desirable level is quite patent. In the last 20 years there have been new theoretical developments and advancements in instrumentation which have helped immensely in the arena of polymeric coatings. The acid-base theory of adhesion has found particular application in controlling the adhesion behavior of coatings. The technical program for this symposium consisted of 23 papers covering many subtopics dealing with adhesion aspects of polymeric coatings. There were lively and illuminating – not exothermic – discussions, both formally and informally, throughout the symposium. The presenters hailed from many corners of the globe and represented varied disciplines and research interests. Now coming to this volume (called Volume 2) it contains a total of 13 papers (others are not included for a variety of reasons) addressing many different issues. It must be recorded that all manuscripts were rigorously peer reviewed and suitably revised (some twice or thrice) before inclusion in this volume. So this book is not a mere collection of unreviewed papers – which is commonly the case with many symposia proceedings – rather it represents information which has passed peer scrutiny. Also it should be pointed out that, for a combination of reasons, the publication of this book got delayed but the authors were asked and given the opportunity to update their manuscripts. So the information contained in this book should be current and fresh.

viii

Preface

The topics covered in this volume include: factors influencing adhesion of polymeric coatings; ways to improve adhesion; formation and relevance of interphase in practical adhesion; adhesion/cohesion in painted plastics; imaging of polymer surfaces; effect of substrate residue (smut) on coating process; surface treatment of metals and glass by silanes; surface modification of polyphenylene sulfide plastics; resin bonding in dentistry; measurement of internal stresses in polymeric coatings; effect of steel surface composition on adhesion of paint; wet adhesion of coatings on wood; and modified tape test to measure adhesion of coatings. Yours truly sincerely hopes that this book will be of interest to everyone interested or involved in the arena of polymeric coatings. Also it should provide some new ideas as to how to control adhesion durability of coatings in different environments. Acknowledgements First, my sincere thanks are extended to my colleague and friend, Dr. Robert H. Lacombe, for taking care of the organizational aspects of this symposium. The comments from the peers are a sine qua non to maintain the highest standard of a publication, so I am most appreciative of the time and efforts of the unsung heroes (reviewers) in providing many valuable comments. I am deeply thankful to the authors for their interest, enthusiasm, patience and contribution without which this book would not have seen the light of day. In closing, my thanks go to the staff of VSP (publisher) for giving this book a body form. K.L. Mittal P.O. Box 1280 Hopewell Jct., NY 12533 1. K.L. Mittal (Ed.), Adhesion Aspects of Polymeric Coatings, Plenum Press, New York (1983).

Adhesion Aspects of Polymeric Coatings, Vol. 2, pp. 1–20 Ed. K.L. Mittal © VSP 2003

Interphase: Formation, characterization and relevance to practical adhesion A.A. ROCHE,∗ J. BOUCHET and S. BENTADJINE INSA de Lyon, Ingénierie des Matériaux Polymères / Laboratoire des Matériaux Macromoléculaires (CNRS, UMR 5627), 20 Ave. Albert Einstein, F-69621 Villeurbanne Cedex, France

Abstract—Epoxy-diamine networks are extensively used as adhesives or paints in many industrial applications. When the precursors are applied onto metallic substrates and cured, an interphase, having chemical, physical and mechanical properties quite different from that of bulk polymer, is created between the substrate and the polymer. Moreover, chemical reactions between diamine and metallic surfaces induce an increase in the practical adhesion (adherence). When the same epoxydiamine mixtures are applied onto gold coated or polyethylene substrates, the interphase properties are the same as bulk ones. When epoxy-diamine mixtures are applied onto aluminum or titanium alloy surfaces, the glass transition temperature, amine and epoxy reaction extent, the interphase thickness, residual stresses within the interphase, Young’s modulus of the interphase all depend on the amine nature (aromatic, aliphatic or cycloaliphatic), the stoichiometric ratio, the processing conditions (time and temperature), the organic layer thickness and the metallic surface treatment. Coating analyses (FTIR, FTNIR, DSC, DMTA, H+ and C13 NMR, SEC, ICP and POM) suggest that diamine monomers chemically react with and dissolve in the metallic hydrated oxide layer. Then, metallic ions diffuse through the organic layer to form a complex by coordination with diamine monomers (chelate or ligand). Metal-diamine complexes are insoluble, at room temperature, both in diamine as well as in DGEBA monomers and they induce a phase separation during the curing cycle of the epoxy-diamine precursors. Furthermore, the chemical bonding of diamine monomers to the metallic surfaces and the diamine-metal crystal orientation parallel to the metallic surface within the interphase lead to chemical, physical and mechanical properties to the epoxy-diamine network which are different from those of the bulk. Keywords: Epoxy-diamine networks; practical adhesion; interphase; diamine-metal chemical reactions.

1. INTRODUCTION

Epoxy-diamine mixtures are extensively used as adhesives or paints in many industrial applications. When they are applied onto metallic substrates and cured, epoxy-amine liquid monomers react with the metallic oxide and/or hydroxide to form chemical bonds [1, 2] increasing practical adhesion (or adherence) between ∗

To whom all correspondence should be addressed. Phone: (33) 4 72 43 82 78, Fax: (33) 4 72 43 85 27, E-mail: [email protected]

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the epoxy polymer and the substrate surface [3, 4]. Different studies concern the influence of the metallic substrate nature on the prepolymer cross-linking. Some authors have shown a catalytic effect of the substrate on the cross-linking mechanisms and the prepolymer adsorption onto metallic surfaces [5]. Others have attempted to determine the existence of monomer/substrate specific reactions. As an example, Dillingham and Boerio [6] have studied the polymerization of a diglycidyl ether of bisphenol A (DGEBA) and a triethylene tetramine (DGEBATETA) system applied onto aluminum by using both FTIR and XPS. Close to the polymer/metal interface, the hardener is partially protonated by aluminum hydroxides. Moreover, the polymerization is catalyzed by the presence of hydroxide acid groups leading to an interphase formation. Nigro and Ishida [7] have studied, by using FTIR, an epoxy/BF3-monoethylamine system applied onto polished steel. The epoxy conversion rate is more important in the polymer/metal interfacial region suggesting the existence of chemical reactions with the steel surface. The epoxy prepolymer was applied onto the steel surface (without hardener), and the same phenomenon was still observed suggesting that the chemical species formed at the steel surface were able to catalyze the homopolymerization of the epoxy monomer. Gaillard and coworkers [8-9] have shown that the DGEBADDA polymerization rate was faster on the zinc surface of galvanized or electrogalvanized steel than on polished steel. They concluded that this was due to a catalytic effect of zinc metallic ions leading to a higher extent of cross-linking. Some researchers have studied the adsorption of various monomers, used as adhesives, at the substrate surface [10-12]. Kollek [13] studied, using FTIR, the polymerization of a diglycidyl ether of bisphenol A and a dicyandiamide (DGEBADDA) system applied onto aluminum and observed that both the DDA and the DGEBA monomers were adsorbed onto the aluminum oxide surface. For the DDA monomer, the adsorption was due to the acidic proton of the aluminum oxide. For the epoxy monomer, the adsorption was achieved by the oxirane (or epoxy) ring opening. Unfortunately, only a few papers [14-17] have dealt with molecular structures formed within the interphase region. Moreover, when epoxy resins are applied onto metallic substrates and cured, intrinsic and thermal residual stresses develop within the entire organic layer [18]. Intrinsic stresses are produced as a result of the mismatch between the active sites of the metallic substrate and the organic network and/or the formation of the polymer network. Thermal stresses are mostly developed during cooling [19] and are the result of thermal expansion mismatch between the metallic substrate and the polymer or cure-induced shrinkage of the organic layer [20]. Whatever their source, these residual stresses reduce the practical adhesion and may induce cracks in coating materials [21-23] resulting in a drop of the overall performance of adhesives or paints. To gain a better understanding of epoxy/metal adhesion requires a full knowledge of chemical and physical reactions which take place in the epoxy/metal interphase [24, 25]. Thus, the polymer/substrate interphase is a complex region containing gradients of residual stresses, as well as allowing structural rearrangement, intermolecular and inter-atomic interactions and diffu-

Interphase: Formation, characterization and relevance to practical adhesion

3

sion phenomenon [24]. When the adhesion of epoxy/metal systems failed, it was possible not only to correlate the residual stresses at the interphase/metal interface to practical adhesion but also to correlate the theoretical adhesion and durability to the presence or not of some chemical species [26]. However, in such systems, chemical and physical mechanisms leading to the interphase formation were not well understood. The aim of this paper was to develop an understanding of this interphase formation and to characterize it using model epoxy-amine systems. 2. EXPERIMENTAL

2.1. Materials 2.1.1. Substrates The metallic substrates used were commercial 0.516 ± 0.005 mm thick 5754 aluminum alloy from Péchiney and 0.600 ± 0.005 mm thick Ti6Al4V titanium alloy from Aérospatiale. Titanium and aluminum sheets were made into squares of dimensions 100 x 100 mm2. Before any monomer application, aluminum and titanium substrate surfaces were ultrasonically degreased in acetone for 10 min. and wiped dry. However, some titanium and aluminum panels were also degreased and chemically treated. Titanium sample were immersed in a solution of 10 g/l of ammonium difluoride for 2 min. at room temperature, rinsed in running tap water for 1 min., immersed in deionized water for 5 min. and wiped dry. Aluminum panels were degreased and submerged in a solution of 250 g/l of sulfuric acid, 50 g/l chromium acid and 44 g/l aluminum sulfate at 60°C for 20 min., rinsed and dried as titanium. After surface treatment, all substrates were kept in an airconditioned room (22 ± 2°C and 55 ± 5% R.H) for 2 h. Some aluminum sheets were coated with gold (≈ 100 nm) using a SCD005 Sputter Coater from Bal-Tec. 2.1.2. Monomers and polymers The bifunctional epoxy prepolymer used was a liquid diglycidyl ether of bisphenol A (DGEBA, M = 348 g/mole, DER 332) from Dow Chemical. The cycloaliphatic diamine curing agent used was isophorone-diamine (IPDA or 3aminomethyl-3,5,5-trimethylcyclohexylamine) from Fluka. An aromatic diamine, (4,4'-methylenebis(3-chloro-2,6-diethylaniline) or MCDEA) from Aldrich and an aliphatic diamine, (polyoxypropylene diamine or D400) from Texaco were also used. Epoxy prepolymer and curing agents were used without further purification. Assuming a functionality of 4 for the diamines and 2 for the epoxy monomer, a stoichiometric ratio (a/e = aminohydrogen/epoxy) equal to 1 was used throughout the work. Homogeneous mixtures of DGEBA and diamine were achieved by stirring under vacuum (≈ 1 Pa) at room temperature for 1 h (Rotavapor RE211 from Büchi, Switzerland) to avoid air bubble formation. The epoxy-amine adhesive cure cycle [27-29] was adapted to obtain both the maximum of the cure conversion, i.e. the highest glass transition temperature denoted Tg∞, and the thickest interphase formation. For IPDA, the curing cycle was: 3 h at 20°C, 20 → 60°C

A.A. Roche et al.

4

(2°C/min.), 2 h at 60°C, 60 → 140°C (2°C/min.), 1 h at 140°C, 140 → 190°C (2°C/min.), 6 h at 190°C, cooling (8 h) in the oven to 20°C; for MCDEA: 3 h at 20°C, 20 → 60°C (2°C/min.), 2 h at 60°C, 60 → 160°C (2°C/min.), 4 h at 160°C, 160 → 190°C (2°C/min.), 9 h at 190°C, cooling (8 h) in the oven to 20°C; and for D400: 3 h at 20°C, 20 → 60°C (2°C/min.), 2 h at 60°C, 60 → 100°C (2°C/min.), 3 h at 100°C, cooling (8 h) in the oven to 20°C. 2.1.3. Sample preparation Several layers of an adhesive tape (5413 from 3M), approximately 50 µm thick, were applied all around the periphery of 100 x 100 mm2 treated metallic sheets to obtain the desired liquid coating thickness. The epoxy-diamine mixtures were poured onto the metallic surfaces and spread with a cylindrical glass rod. For bulk materials, 10 x 10 x 100 mm3 bars were prepared using a silicone (RTV 501 from Rhône-Poulenc) or PTFE mould. Only the central parts of the bars were used for analysis. After curing and cooling down, the coating thicknesses (from 40 to 1500 µm) were determined using an EG-100 Digital Linear Gauge (from Ono Sokki Co, Japan) having a ± 2 µm sensitivity. 2.1.4. Monomer analysis To characterize the changes in monomers, both the liquid monomers DGEBA and IPDA were applied between two chemically treated metallic substrates (100 x 50 mm2) to form a 110-150 µm thick liquid film and kept at room temperature for 3 h in a desiccator under continuous nitrogen flow to prevent any monomer carbonization or oxidation. Then, the liquid monomers were scraped from the metallic surfaces (and will be called “modified monomers” in the following) with a PTFE spatula and stored in polyethylene vials under a nitrogen atmosphere. 2.2. Experimental techniques 2.2.1. Differential scanning calorimetry (DSC) DSC experiments were carried out in a Mettler DSC 30 apparatus to determine the glass transition temperature (Tg) of epoxy resins. Sealed aluminum pans containing 15-20 mg of resin were heated from –50°C to 250°C at a rate of 10°C/min under a continuous flow of U-grade argon. Samples were weighed using a Mettler balance having a ± 1 µg sensitivity. The calorimeter was calibrated with both indium and zinc. The glass transition temperature (onset) was determined with a ± 0.5°C sensitivity. To evaluate the variation in Tg with the coating thickness, the (Tg)eq value at thickness (i) was calculated using: (Tg ) eq =

hiTgi − hi−1Tgi−1 hi − hi−1

(1)

where Tgi corresponds to the Tg of a hi thick coating and Tgi–1 value corresponds to the glass transition temperature of a hi–1, thick coating.


5

2.2.2. Dynamic mechanical thermal analysis (DMTA) Dynamic viscoelastic measurements on coatings, after removal from the metallic substrate, were performed in a Rheometrics Solids Analyzer RSA II apparatus, using the tensile mode to determine the storage (E'), and loss (E") moduli as well as the ratio E"/E' = Tan δ as a function of temperature. Samples (4x40 mm2) were heated from –150°C to 250°C using a heating rate of 2°C/min by a forced convection oven using a nitrogen stream. The sample was deformed sinusoidally to a controlled strain amplitude of 0.05% at a fixed frequency of 1 Hz. Measurements were made at 30 s intervals. 2.2.3. Fourier transform infrared and near-infrared spectroscopy (FTIR, FTNIR) An infrared spectrometer (FTIR Magna-IR 550 from Nicolet) was used with Omnic FTIR software. An Ever-GloTM source was used along with a KBr beamsplitter and DTGS-KBr or MCT/A detector. The mid-infrared spectra were recorded in the 400-4000 cm–1 range and in the 4000-7000 cm–1 range for the nearinfrared. A transmission accessory was used for bulk or free standing film characterization. For the bulk material, KBr was mixed in a 100:1 ratio with the various epoxy resins which had been ground cryogenically. This mixture was then pressed under vacuum to obtain disks. For analysis, pure KBr disks were used as background. For each analysis, 64 or 96 scans were collected at 4 cm–1 resolution. The cure conversions for epoxy and amine groups were calculated by using, respectively, the 4530 cm–1 epoxy combination band and the 6500 cm–1 amine band [30]. The 4623 cm–1 aromatic C-H ring stretch combination band was considered as reference. Thus, the amine (Xa) and the epoxy (Xe) conversions were determined using the ratio of the respective peak areas (A) by: Xa =1 −

( A6500 / A4623 )t ( A6500 / A4623 )t=0

and

Xe =1 −

( A4530 / A4623 )t ( A4530 / A4623 )t=0

(2)

Once again, to evaluate the Xa and Xe variation versus the coating thickness we calculated the (Xa)eq and (Xe)eq values at thickness (i) using:

( X a,e )eq =

(

hi X a,e

)i − hi−1 ( X a,e )i−1 hi − hi−1

(3)

According to Bentadjine [31], it is possible to correlate the stoichiometric ratio (a/e) to the Xa/Xe ratio as shown in Fig. 1. Then, the variation of the equivalent stoichiometric ratio (a/e)eq versus the coating thickness can be determined. 2.2.4. Nuclear magnetic resonance (NMR) spectroscopy For proton and carbon Nuclear Magnetic Resonance (1H and 13C NMR) a Bruker AC 400 spectrometer was used with the sodium salt of deuterated trimethyl silyl propionic acid (TSPd4) as the internal standard. Analyses were carried out at 293 K using deuterated water in which pure IPDA and modified IPDA were com-

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Figure 1. Variation of the (Xa)eq/(Xe)eq ratio as a function of the stoichiometric ratio (a/e).

pletely soluble. By comparing pure IPDA and modified IPDA spectra it was possible to determine the differences in chemical shifts for the equivalent carbon nuclei. 2.2.5. Inductively coupled plasma (ICP) spectroscopy An ICP spectrometer (Modula by Spectro Analytical Instruments) was used with a 2.5 kW plasma generator at 27 MHz. It was fitted with various detectors (an UV (0.75 m, 2400 grooves mm–1; 160-480 nm) monochromator, UV (0.75 m, 3600 grooves mm–1) polychromator, and visible (0.75 m, 1200 grooves mm–1) polychromator). A cross-flow nebulizer was used to introduce the liquid sample. Distilled water was used for dilution. 2.2.6. Size exclusion chromatography (SEC) A Waters apparatus was used with double detection (UV Waters 484 Tunable Absorbance Detector at λ = 254 nm and Waters 410 differential Refractometer) and a Waters 510 HPLC pump. The elution solvent used was tetrahydrofuran (THF) and the separation was carried out on two styrene-divinyl benzene columns (Nucleogel 100-5 and 500-5 from Macherey-Nagel) with a flow rate of 1 ml/min. The average molar mass calibration curve was constructed from monodispersed polystyrene standards and used to correlate the retention time to the mass average molar mass (M) .


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2.2.7. Polarized optical microscopy (POM), scanning electron microscopy (SEM) and electron microprobe analysis (EMPA) Drops of diamine or epoxy-diamine mixture were confined between two glass plates and mounted on a hot plate under the optical microscope (Laborlux 12POLS from Leica equipped with a hot plate FP82 from Mettler and a CCD-IRIS color video camera from Sony). Samples were heated from 30°C to 190°C at 10°C/min. A scanning electron microscope (Philips XL20) fitted with an electron microprobe analysis accessory (Edax-Econ4) was also used. Samples were neither coated with gold nor with carbon. The accelerating electron voltage was 5 kV, the electron spot diameter for microanalyses was about 200 nm and the tilt angle used was 15°. 2.2.8. Free corrosion potential and pH measurements Different aqueous solutions with isophorone diamine (IPDA) were prepared. The pH’s of the solutions were measured at room temperature (22°C) using a PHN 850 pH-meter from Tacussel equipped with a combined electrode (Pt-Ag/AgCl). The pH variation versus the diamine concentration in aqueous solution is shown in Fig. 2. The basic behavior of IPDA is clearly shown. For DGEBA/IPDA prepolymer with a/e = 1, the IPDA concentration is equal to 20 wt% which corresponds to a pH value of 12. The same basic behavior was obtained for aqueous D400 (pH = 12.1) and MCDEA (pH = 10.3) monomer solutions. Free corrosion potential (or open circuit potential) measurements were carried out by using a Tacussel Potentiostat (PRT 40-1X) with a combined Pt electrode (Ag/AgCl).

Figure 2. pH variation of aqueous solution as a function of IPDA concentration.

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3. RESULTS AND DISCUSSION

3.1. Bulk and coating characteristics According to the curing cycle mentioned above, the maximum glass transition temperatures (Tg∞) of DGEBA/IPDA, DGEBA/MCDEA and DGEBA/D400 bulk epoxies were 163°C, 188°C and 34°C respectively. For these bulk materials, the amine conversion (Xa) and the epoxy conversion (Xe) were equal to 1. These values are in good agreement with the previous works [27-29]. Variations in the glass transition temperatures (Tg)eq versus DGEBA/IPD coating thickness applied onto both degreased titanium and aluminum are shown in Fig. 3. Variations in the equivalent stoichiometric ratio ((a/e)eq) versus the thickness of DGEBA/IPDA coatings applied onto degreased titanium and aluminum substrates are shown in Fig. 4. Regardless of coating thickness, for the coatings applied onto a gold coated substrate, the (Tg)eq’s were equal to the (Tg∞) of the cured bulk DGEBA/IPDA material and the equivalent stoichiometric ratio ((a/e)eq) was equal to 1. Similarly, for coatings applied either onto degreased titanium or aluminum that were thicker than 400 µm or 200 µm respectively, (Tg)eq and (a/e)eq were all equal to those of the cured bulk DGEBA/IPDA system. For the thinnest coatings (≤ 50 µm) applied onto titanium or aluminum, (Tg)eq values were quite different (≈ 129°C and 108°C respectively) from the bulk ones (163°C) and the equivalent stoichiometric ratios (a/e)eq were 1.19 and 1.14, respectively. Thus it can be assumed that for these thin coatings, a different epoxy network was formed. For coating thicknesses between 50 ≤ h ≤ 400 µm for Ti and 200 µm for Al, a gradient region was observed. In all the following work, the interphase thickness (or width) will be defined as the region where the properties are different from those of the bulk. In the case of the DGEBA/IPDA coatings, the interphase is thicker for the titanium substrate (400 µm) than for the aluminum one (200 µm). So, we can assume that the dissolution rate of TiO2 amorphous oxide is higher than that of the Al2O3 amorphous oxide. According to Ellingham [32], the standard free energies at room temperature of formation of Al2O3 and TiO2 crystalline oxides are equal to 250 and 210 kcal/mol O2, respectively. In a first approximation, we can assume that the amorphous oxide layer formed on aluminum alloy was more stable than the one obtained on titanium alloy. Thus, for the same diamine monomer, the dissolution rate of TiO2 should be higher than that of Al2O3 leading to a wider interphase for TiO2, and this is observed experimentally. The large interphase thickness also means that for a 1 mm thick coating on degreased titanium, the first 400 µm from the substrate surface corresponds to the interphase and that only the remaining 600 µm will have the same properties as the DGEBA/IPDA bulk epoxy. Also, we observe that the interphase thickness depends on the substrate nature (Fig. 4) as well as on the basic behavior of the amine as shown in Fig. 5. The formation of such a thick interphase has to be explained.


9

Figure 3. Variation of the equivalent glass transition temperature versus the coating thickness for degreased titanium and aluminum substrates.

Figure 4. Variation of the equivalent stoichiometric ratio (a/e)eq as a function of the coating thickness for both degreased titanium and aluminum substrates.

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Figure 5. Variation of the interphase thickness versus the pH of monomer aqueous solution.

3.2. Monomers characterization To characterize the chemical changes occurring, DGEBA and IPDA liquid monomers were placed between two chemically treated metallic substrates (Ti or Al) and kept at room temperature for 3 h. Then, Ti or Al modified monomers were scraped from their metallic substrates with a PTFE spatula and visually compared to the pure monomers. Both pure monomers (DGEBA and IPDA) were transparent. No color change of the DGEBA monomers after application onto titanium or aluminum was observed. The IPDA monomers from both aluminum and titanium became milky white (Fig. 6). After two hours within test tubes, a separation was observed. The floating part remained transparent while the precipitate was milky white. One drop each of the precipitate from both Ti and Al modified IPDA were placed between two glass plates and observed using a polarized optical microscope (Fig. 7). Crystals in the modified IPDA liquid were observed corresponding to a new chemical product formed when the IPDA monomer reacted with the metallic oxide and/or hydroxide surface of both titanium and aluminum. Taking into account the basic nature of the pure diamine curing agent, a dissolution of the metallic oxide or hydroxide by the amine is assumed when the monomer is applied onto the metallic substrate. The free potential variation versus the immersion time for chemically etched titanium immersed in a 20 wt% aque-


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Figure 6. Photographs of pure IPDA and IPDA modified either on titanium or aluminum.

Figure 7. POM micrographs of the modified IPDA precipitates obtained after application onto titanium (a) and aluminum (b) substrates.

ous IPDA solution is reported in Fig. 8 and it suggests a dissolution of the metallic oxide surface. A comparison of the SEM micrographs (Fig. 9) of the degreased titanium and aluminum surfaces to the ones obtained after an application of IPDA for 3 h and rinsing with acetone shows indisputably a chemical attack of the metallic surfaces by the liquid IPDA monomer. The dissolution of metallic oxide or hydroxide surface layers may induce a metallic ion diffusion within the IPDA liquid monomer. Experimentally, this was observed by ICP analysis of pure IPDA and the precipitate portion of Ti and Al modified IPDA as reported in Fig. 10. The presence of titanium or aluminum is shown clearly in the modified IPDA monomer spectrum. Thus the hypothesis of the dissolution of the titanium or aluminum oxide and/or hydroxide by the basic diamine monomer is strengthened. FTIR spectra of pure IPDA and modified IPDA after application onto aluminum,

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Figure 8. Variation of the free potential versus time for a chemically etched titanium substrate immersed in an aqueous solution of IPDA (20 wt%).

Figure 9. SEM photomicrographs of the initial surfaces of degreased titanium and aluminum substrates and after application for 3 h of liquid IPDA monomer and removal using acetone rinse.


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Figure 10. ICP spectra of pure IPDA and titanium and aluminum modified IPDA.

Figure 11. FTIR spectra of pure IPDA and IPDA modified on titanium, aluminum or gold coated substrates.

titanium or gold coated substrate are shown in Fig. 11. The spectra of pure IPDA and modified IPDA after application onto a gold coated substrate or the floating portion of the modified IPDA on aluminum and titanium are identical. However, new bands (between 1100-1350 cm–1 and 600-700 cm–1) are observed only when IPDA was applied onto titanium and aluminum substrates. FTIR results confirmed that no chemical reaction occurred when coatings were applied onto the gold coated substrates. The diamine monomer (IPDA) clearly reacts with both titanium and aluminum substrates. Fig. 12 shows the reflectance index (RI) SEC spectra obtained from pure diamine monomer (MCDEA) and modified MCDEA with both titanium and aluminum substrates. New small peaks are observed which correspond to M w = 1014 g/mol for titanium and M w = 1011 g/mol for aluminum. These values are twice higher than the pure MCDEA monomer mass average molar mass ( M w = 489 g/mol).

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Figure 12. SEC spectra of pure MCDEA and MCDEA modified on titanium and aluminum substrates.

To characterize the chemical bindings that may appear between the IPDA monomer and metallic ions, nuclear magnetic resonance (NMR) spectroscopy was used. Pure IPDA monomer was actually a mixture of 25% “cis” and 75% “trans” isomers and the corresponding conformations and the carbon atom labeling are shown in Fig. 13. Peak assignment of pure IPDA was achieved using 1H, 13 C and 2D NMR spectra as reported elsewhere [33, 34]. Spectra of pure IPDA monomer and modified IPDA applied onto titanium and aluminum were also obtained. After a complete peak assignment, it can be noted that some of the characteristic nuclear resonance frequencies vary (magnetic shielding) when IPDA was applied onto titanium and aluminum. From the differences between characteristic resonance frequencies, chemical shifts can be calculated [35]. Fig. 13 shows also the NMR chemical shifts for various carbon atoms when IPDA was applied onto titanium and aluminum. Only carbon atoms associated with nitrogen atoms or neighboring carbon atoms were affected when the monomer was applied onto metallic substrates and the shift was found to be the same regardless of the isomer. The chemical shifts obtained for aluminum were about two to three times higher than the ones observed on titanium. From these results, it can be assumed that the bond between amine groups and metallic ions results from the electron-rich nitrogen atom donating its lone pair to the electron-deficient metallic center. Thus, organo-metallic complexes were created by coordination bindings. When the com-


15

Figure 13. Cis and trans isomers configurations and labeling of the two IPDA isomers and chemical shifts observed for modified IPDA on titanium and aluminum.

plex concentration within the liquid monomer reaches the solubility limit (Ks), complexes crystallize as sharp needles as shown in Fig. 7. About forty years ago, the complex formation was reported for organic pigments where metallic ions were introduced to stabilize monomers [36]. More recently other authors have described the complex formation between amines and metals [37-38]. Using such modified monomers, it was interesting to add them to the DGEBA epoxy prepolymer and to characterize the properties of such new materials. 3.3. Pure DGEBA/modified IPDA characterization To fully understand the properties of interphases when DGEBA-diamine mixtures are applied onto metallic surfaces and to verify our various hypotheses, we prepared some epoxy mixtures with pure DGEBA and both Ti and Al modified IPDA monomers. The stoichiometric ratio was kept equal to 1 and the curing conditions were the same as for the pure DGEBA/IPDA system. Figure 14 contains POM photographs obtained before and after the cure cycle for both pure DGEBA/titanium or aluminum modified IPDA mixtures. Comparing Fig. 7 to Fig. 14, a dilution effect was observed before curing, but crystals were still present in the pure DGEBA/Ti or Al modified IPDA mixtures. After curing, crystals remain observable and all around these crystals, a phase separation (black parts in

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Figure 14. POM photographs of mixtures (a/e = 1) of pure DGEBA/Ti and Al modified IPDA.

the photograph) can be observed. It is quite evident that the presence of crystals and the phase separation lead to the formation of a new network. The properties of such modified polymers are reported in Table 1 and are compared to thin coatings on titanium and aluminum and to the pure bulk polymer. The properties of the pure DGEBA/Ti or Al modified IPDA are the same as those of thin films formed onto titanium or aluminum substrates. This means that we have been able to reproduce within a modified bulk system the various phenomena observed for the thin coatings. Moreover, an increase of Young’s modulus of DGEBA/modified IPDA systems was obtained as shown in Fig. 15. Randomly dispersed organo-metallic crystals act as short fibers in a polymer matrix [39-43]. However, this 50% increase in Young’s modulus does not explain the huge increase (300%) in the longitudinal Young’s modulus of thin films as shown in Table 1. Lastly, one drop of IPDA monomer was applied close to an aluminum surface between two glasses. The POM photographs were taken versus time. After


17

Table 1. Physical, chemical and mechanical properties of bulk coating and of thin films applied onto degreased titanium and aluminum substrates

DGEBA/IPDA bulk 40 µm thick film on Al DGEBA/Al modified IPDA 250 µm thick film on Ti DGEBA/Ti modified IPDA

Tg [°C]

a/e

E [GPa]

163 110 108 129 132

1.0 1.14 1.15 1.20 1.22

3.2 10 5.0 11 4.8

E: is the Young modulus a/e: is the stoichiometric ratio

Figure 15. Tensile stress versus strain curves for pure DGEBA/IPDA system and pure DGEBA/Al modified IPDA system.

30 min, the first crystals appeared and after 3 h it can be clearly observed that most of crystals are oriented parallel to the metallic surface as shown in Fig. 16. Thus, the huge increase in the longitudinal Young’s modulus as determined by the flexure test can be easily explained.

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Figure 16. POM photographs of liquid IPDA monomer in the vicinity of degreased aluminum surface after 30 min (left) and 3 h (right).

4. SUMMARY

When epoxy-diamine prepolymers are applied onto metallic substrates, an interphase between the coating part having the bulk properties and the metallic surface was created. Chemical, physical and mechanical properties of the interphase formed depend on the substrate nature and its surface treatment, the nature of the diamine hardener, and the curing cycle. Use of an appropriate curing cycle for the prepolymer allowed us to reach the maximum cure conversion of the polymer and the full interphase formation. Mechanisms describing the formation of the interphase were deduced from comparison of behaviors when epoxy and diamine monomers were applied onto aluminum and titanium. When either pure DGEBA or diamine monomers were applied onto gold coated surfaces, no chemical reaction was observed. In the same way, when the pure DGEBA monomer was applied onto the metallic surfaces, no chemical reaction was observed. On the contrary, when pure diamine monomers were applied onto metallic surfaces, chemical reactions occurred. According to the basic behavior of diamine monomers, a partial dissolution of the surface metal oxide and/or hydroxide was observed. Metallic ions diffuse within the liquid monomer mixture (epoxy-diamine) and react with the amine groups of the diamine monomers to form organometallic complexes by coordination. When the complex concentration is higher than the solubility product, these complexes crystallize as sharp needles. They align themselves parallel to the metallic surfaces leading to an oriented crystalline layer in the vicinity of the substrate surface. Crystals act as short fibers in an organic matrix leading to an increase of the mechanical properties. During the curing cycle, crystals were not fully dissolved and a phase separation was observed to occur all around these crystals inducing the formation of a new epoxy network. The same phenomena would be expected for all metallic substrates as soon as they are covered with an oxide or hydroxide layer. Also, since dissolution and diffusion phenomena are expected, the extent of interphase formation should be related to the contact duration between the liquid diamine and metallic substrates.


19

This enabled us to propose interphase formation mechanisms leading to a better understanding of its chemical, physical and mechanical properties. Acknowledgements We are very grateful to R. Petiaud and V. Massardier for performing the NMR analyses and to R. Diemiaszonek for the ICP analyses. REFERENCES 1. 2. 3. 4. 5. 6. 7. 8.

E. Peillex, Ph.D. thesis, Université de Lyon 1 (1996). C. Fauquet, Ph.D. thesis, Université Paris 6 (1992). P. Walker, J. Coating. Technol., 52, 49 (1980). R.D. Guminski and F.M.P. Meredith, J. Oil. Colour Chem. Assoc., 44, 93 (1961). B. De’neve, Ph.D.thesis, Ecole Nationale Supérieure des Mines de Paris (1993). R.G. Dillingham and F.J. Boerio, J. Adhesion, 24, 315 (1987). J. Nigro and H. Ishida, J. Appl. Polym. Sci., 38, 2191 (1989). F. Gaillard, H. Hocquaux, M. Romand and D. Verchère, Proceedings of Euradh’92, pp. 122127, Karlsruhe (Germany) (1992). 9. D. Verchère, H. Hocquaux, T. Marquais and F. Gaillard, Proceedings of Euradh’92, pp. 488495, Karlsruhe (Germany) (1992). 10. J.E. De Vries, L.P. Haack, J.W. Holubka and R.A. Dickie, J. Adhesion Sci. Technol., 3, 203 (1989). 11. S.G. Hong, N.G. Cave and F.J. Boerio, J. Adhesion, 36, 265 (1992). 12. J.W. Holubka and J.C. Ball, Ind. Eng. Chem. Res., 28, 48 (1989). 13. H. Kollek, Int. J. Adhesion Adhesives, 5, 75 (1985). 14. J. March, L. Minel, M.G. Barthes-Labrousse and D. Gorse, Appl. Surface Sci., 133, 270 (1998). 15. G.S. Crompton, J. Mater. Sci., 24, 1575 (1989). 16. Y.H. Kim, G.F. Walker, J. Kim and J. Park, J. Adhesion Sci. Technol., 1, 331 (1987). 17. D.L. Allara and C.W. White, Amer. Chem. Soc., 273 (1978). 18. A. Entenberg, V. Lindberg, L. Fendrock, Sang-Ki Hong, T.S. Chen and R.S Horwath, in: Metallized Plastics 1: Fundamental and Applied Aspects, K.L. Mittal and J.R. Susko (Eds.), pp. 103113, Plenum Press, New York (1989). 19. F.J. Von Preissing, J. Appl. Phys., 66, 4262 (1989). 20. P. Scafidi and M. Ignat, J. Adhesion Sci. Technol., 12, 1219 (1998). 21. H. Orsini and F. Schmit, J. Adhesion, 43, 55 (1993). 22. M.D. Thouless and H.M. Jensen, J. Adhesion Sci. Technol. 8, 579 (1994). 23. D.R. Mulville and R.N. Vaishnav, J. Adhesion, 7, 215 (1975). 24. K.L. Mittal, in: Adhesion Measurement of Thin Films, Thick Films and Bulk Coatings, K.L. Mittal (Ed.), pp. 5-17, ASTM, STP 640, Philadelphia (1978). 25. L.H. Sharpe, J. Adhesion, 4, 51 (1972). 26. M.F. Finlayson and B.A. Shah, J. Adhesion Sci. Technol., 4, 431 (1990). 27. O. Sindt, J. Perez and J.F. Gerard, Polymer, 37, 2989 (1995). 28. L. Bonnaud, Ph.D. thesis, INSA de Lyon (1999). 29. H. Masood Siddiqi, Ph.D. thesis, INSA de Lyon (1997). 30. N. Poisson, G. Lachenal and H. Sautereau, Vibrational Spectros., 12, 237 (1996). 31. S. Bentadjine, Ph.D. thesis, INSA de Lyon (2000). 32. H.J.T. Ellingham, J. Soc. Chem. Ind., 63, 125 (1944). 33. S. Bentadjine, A.A. Roche and J. Bouchet, in: Adhesion Aspects of Thin Films, K.L. Mittal (Ed.), Vol. 1, pp. 239-260, VSP, Utrecht (2001). 34. S. Bentadjine, R. Petiaud, V. Massardier and A.A. Roche, Polymer, 42, 6271-6282 (2001).

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35. F.K. Bovey, High Resolution NMR of Macromolecules, Academic Press, London (1972). 36. S. Wernick and R. Pinner, Les traitements de surface et la finition de l’aluminium et de ses alliages, Eyrolles, Paris (1962). 37. M. Kinzler, M. Grunge, N. Blank, H. Shenkel and I. Scheffler, J. Vac. Sci. Technol., 10, 26912697 (1992). 38. J.S. Kuo and J.W. Rogers, Surface Sci., 453, 119-129 (2000). 39. J. Bouchet, A.A. Roche, E. Jacquelin and G.W. Scherer, in: Adhesion Aspects of Thin Films, K.L. Mittal (Ed.), Vol. 1, pp. 217-237, VSP, Utrecht (2001). 40. M.G. Bader and W.H. Bowyer, J. Phys. D: Appl. Phys., 5, 2215-2226 (1972). 41. A.R. Saadi and M.R. Piggott, J. Mater. Sci., 20, 431-437 (1985). 42. M. Miwa and N. Horiba, J. Mater. Sci., 29, 973-977 (1994). 43. M.L. Mehan and L.S. Schadler, Composites Sci. Technol., 60, 1013-1026 (2000).


Depletion, a key factor in polymer adhesion G. FRENS∗ Laboratory of Physical Chemistry, Delft University of Technology, Julianalaan 136, 2628 BL Delft, The Netherlands

Abstract—The chemical properties of polymers are essential for the strength of an adhesion bond. Generally, a random coil of a polymer is repelled, entropically, from an impenetrable surface by the depletion effect. To bond the polymer molecules it must be forced actively towards the interface. Suppression of depletion is brought about by adsorption of specific groups, interacting with polar sites on the substrate. Altering the balance between adsorption and depletion is the most important effect of pretreatments on the strength of an adhesion bond. Adsorption diminishes the effective distance between the surface and the polymer. This, rather than the effect of polar groups on the work of adhesion, leads to order of magnitude increases of the adhesion force between a polymer and a surface. Keywords: Depletion; adhesion; polymers; polar groups; adsorption.

1. INTRODUCTION

A starting point for quantitative adhesion theories is the work of adhesion WA, expressed as: WA = γS + γL – γSL

(1)

The energy WA, of the adhesion between a solid substrate S and a wetting liquid L, follows from the surface energies (in mJ/m2) γS and γL and from their interfacial energy γSL. WA can be determined from contact angle measurements using groups of standard liquids [1-2]. The idea behind the work of adhesion is that when the solid and the liquid are separated, two free surfaces (of the solid S and of the liquid L) are created out of an equal area of contact which had, originally, the interfacial energy γSL. Polymer molecules which are attached to the substrate surface are instrumental in the strength of an adhesion bond. When forces are applied to separate the bond, there is a concentration of stresses at the interface. The stored mechanical energy is released when a crack develops. Propagation of the crack can be stopped, and the bond is saved, when the stresses are redistributed away from the tip of the ∗

Phone: +31 15 278 5180, Fax: +31 15 278 4135, E-mail: [email protected]

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crack. This is what the relatively large and interwoven polymer coils can do for the strength of an adhesion bond. The polymer molecules can only have such an effect if they are fixed strongly onto the surface. This fixation is the chemical secret in the formulation of adhesives for different materials. Pre-treatment procedures are designed for the same purpose: to optimize the strength of the bond, by attaching the polymers to the bonding substrate. But the mechanism which determines the strength of the bond is not very well understood, as yet, from a theoretical point of view. New observations about the adhesion between polymers and solid surfaces [3] indicate why theories of adhesion which are based on the value of WA [4] have always had serious trouble to adequately predict the strength of the bond between a polymer and the surface of a solid material. This is because the effective distance H between the polymer and the substrate has been overlooked as a variable. It can vary by an order of magnitude, depending on experimental conditions which affect polymer adsorption. 2. ADHESION STRENGTH

Surface energies are the macroscopic expression of the intermolecular (van der Waals) attractions between the molecules near the interface. From the Hamaker equation [5] for the long range dispersion forces across the interface it follows that: VA = WA (δ/H)2,

(2)

VA is the attraction energy between two slabs of materials - such as a solid and an adhesive - separated by a gap of width H. The molecular parameter δ can be understood as the distance of closest approach between the two molecules which are in contact across the interface. The adhesion force, FA is the derivative; i.e., FA = dVA/dH

(3)

FA is the force (per m2) which resists the separation of the materials by outside forces. It represents the strength of the adhesion bond and decreases as the cube of the distance H. The bond strength is, therefore, not only dependent on WA, but also on the effective distance H between the bulk polymer and the surface of the substrate. This may seem trivial when the two materials are in close contact, and, therefore, its consequences are often overlooked in theories of adhesion. These theories then proceed and concentrate on methods to maximize WA for making strong adhesion bonds, e.g. through pre-treatments of the substrate surface which lead to an extra (“polar”) attraction between molecules in close contact at H = δ. Indeed, for a strong bond WA must be large. But it is more important that the effective distance H has the minimum possible value. It is in this aspect of adhesion that the physical chemistry of polymer molecules and of the solvent has a special role to play. Polymer molecules in a solvent or in a melt will keep their distance to the interface at H δ.

Depletion, a key factor in polymer adhesion

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3. POLYMER SOLUTIONS

Adhesives are fairly concentrated polymer solutions. In cured and dried layers of adhesives, like in a hot melt of a polymer, the per volume concentration Cs of polymer segments approaches that of the pure polymer. In a dilute polymer solution the volume of the random coil is characterised by the extension vector R. The value of R is related to the size of the segment A and the number of segments N in the coil by áR2ñ= N6/5 A2. When the polymer concentration increases, coils of this size will fill the complete volume of the solution at a segment concentration Cs*. At higher concentrations than Cs* the polymer coils interpenetrate. Such a system, with a concentration above Cs* is called semi-dilute. The polymer coils form a network with a characteristic length scale ξ, which is called the correlation length. At Cs*, the correlation length is approximately equal to R. In a semi-dilute solution, where Cs > Cs*, the relation between ξ and Cs can be found from a scaling argument. This relation can always be written as a power law, such as ξ = R (Cs* / Cs)x with an unknown power x. The correlation length ξ is a local parameter, which describes the interpenetrating network. As such, it cannot depend on the lengths of the individual coils. From the theory of a dilute polymer solution in a good solvent it follows that R is proportional to N3/5 and that Cs* is proportional to N–4/5. The power x in the relation between ξ and Cs must, therefore, be 3/4. With Cs = nN/V, in which n is the number of polymer coils in the solution, and β as the excluded volume of one polymer segment of length A, it follows for a semi-dilute system (Cs Cs*). ξ = β–1/4 · A–1/2 · Cs–3/4

(4)

The correlation length determines, among others, the modulus of elasticity in the viscoelastic adhesive material [6], and therefore the rupture dynamics in thin viscoelastic films [7]. In Eq. 4 A is the length of the Kuhn segment in the polymer chain. The excluded volume β of the segment describes its interaction with the solvent. At high volume fractions of the polymer, like in a melt, the correlation length ξ decreases. It approaches the segment length A as ξ = A · φ–3/4 in the limit of a pure polymer without solvent. The Kuhn segment in a polymer molecule is always larger than just one single monomer. In flexible polymers the length A is typically of the order of 1 nm. It may, however, sometimes be much larger – with stiff macromolecules or in special solvents. The segment length A, however, is always the lower limit for the correlation length ξ in a polymer solution when its concentration increases. 4. DEPLETION AND ADSORPTION

All polymer molecules tend to avoid contact with an impenetrable surface. This effect is known as “depletion”. The entropy of the polymer coils would decrease, if they should approach a surface to distances (H) smaller than the correlation

G. Frens

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length (ξ). This would raise the free energy of the system, and, therefore, polymer molecules leave a gap of width H = ξ at the interface. The free energy involved in this depletion effect is of the order T∆S = kT/ξ 2

(5)

2

per m of contact surface. The effective distance H for the most intimate contact between a polymer coil and a completely inert, but impenetrable, surface is then at least of the order of ξ ≈ A ≈ 1 nm and depends on the stiffness of the chain of segments. This implies, indeed, that for a concentrated solution in a good solvent or a polymer melt the mininum thickness of the depletion layer H δ. The contribution of polar interactions to the adhesion will remain very small as long as H > δ. The adhesion is then mostly determined by the long range dispersion forces across the interface. For strong adhesion this depletion effect must be suppressed. To bring polymer molecules nearer to the interface, the repulsion force due to depletion has to be compensated by attraction between the polymer segments and the surface material. In adhesive systems this attraction is through the specific energies of interaction between polar adsorption sites on the surface and segments of the adjacent polymer chains. Given the right circumstances, the polar groups may pull the polymer chains nearer to the surface. For this, the adsorption energy of the polar groups in segments which attach themselves to the surface must exceed the repulsive depletion energy of kT/A2 (per m2 of interface). The adsorption which determines the quality of the adhesion bond can be measured in the same way as WA, through the determination of contact angles. The method has been described by Pennings [8-11], who showed that for many polymers: ∆γ/γp = α[γs /γp]1.85

(6)

Eq. 6 is called the Pennings Equation. It is used in polymer technology to establish optimum process conditions for precision injection moulding. In the Pennings equation ∆γ is the difference between the surface tension immediately after the contact with the substrate was disjoined and the equilibrium surface tension γp of the polymer layer in air. It registers the presence of the extra groups which have become attached to the surface by adsorption. For optimum bond strength ∆γ can be influenced by pre-treatment of the substrate surface. Pretreatments create more polar groups at the surface and add an extra “polar” contribution to the surface energy γs. The parameter α is the coil compliance of the specific polymer [3], which is a measure of its entropic resistance to the conformation change when it becomes adsorbed. The coil compliance can be influenced by the choice of the solvent, by the temperature and by other process conditions. These may be adjusted and define a technology window for reaching an optimum value of ∆γ and of the bond strength which depends on the adsorption of polymer segments directly onto the bonded surface.


25

5. THE ATTRACTION RANGE, H

If the depletion is overcome by the adsorption of the polymer molecules, the effective distance decreases from H = ξ to H = δ, i.e. from about 1 nm to atomic dimensions. The contact then is as if the substrate were wetted with a liquid instead of covered with polymer segments. At first sight, the difference in H may seem a subtle change indeed, but the strength of the adhesion varies as H 3. Therefore, suppression of the depletion and filling up the gap H between the polymer and the substrate strengthens the adhesion bond by at least an order of magnitude. This is purely an effect of the van der Waals dispersion forces between the bulk of the two materials. These forces become stronger at close range. As a consequence of the interaction of polar groups in the interface the gap width H is reduced, and this has a much larger effect for the bond strength than whatever small contributions the polar groups themselves may have to WA. It is not the work of adhesion which is an indication for the strength of an adhesion bond, but the distance H (of the order of nm) between the bonded materials being reduced by adsorption. An example of the increased adhesion which is created by a pretreatement which generates polar groups on the surface of a substrate is given [12] in Figure 1. Polypropylene samples were treated at different UV-wavelengths in air. This radiation forms ozone, which is adsorbed and reacts at the surface of the polymer so that polar groups are introduced. The surface energy γS of the material increased from 18 mJ/m2 to 52 mJ/m2. Afterwards the samples were glued together with MS-polymer adhesive (a silylated thermosetting adhesive with a polypropylene-ether backbone). The pretreatment led, compared to cured but untreated sam-

Figure 1. Effect of UV treatment of polypropylene on its adhesion strength.

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G. Frens

ples, to a more than tenfold increase in the adhesion strength when glued together parts were tested for the tensile strength of the adhesive bond. This effect seems too large to be explained from the increase in WA alone. The strongest effect must be due to the reduction of the distance H which is brought about by the adsorption of the MS-polymer on the pre-treated polypropylene substrate. 6. DELAMINATION

For a strong adhesion bond it is essential that the depletion effect is suppressed. The competition between adsorption and depletion will, however, always lead to a preferential adsorption of polar molecules, like water or monomers, at the interface. These small molecules have comparable adsorption energies as the segments, but there is no depletion energy to compete with it. Replacing the adsorbed polar groups of the polymer at the interface with water can undo the adsorption of the polymer segments and activate depletion. Sometimes this is a useful effect, e.g. when release agents separate injection moulded plastic products from their metal mould. But for polymer adhesion bonds this sort of depletion is catastrophic. It causes delamination along the interface between the two materials. Therefore, in order to prevent delamination, the adsorption energy of the polar groups on the substrate surface must always exceed that of water by the (depletion) energy kT/A2. The tendency towards depletion is always lurking, hidden in every polymer adhesion bond which has, of necessity, polymer molecules in the interface. This is the reason for interfacial failure in the presence of small polar molecules like monomers or H2O. Such molecules compete with the polymer segments for the polar sites at the surface. Adsorption of water molecules instead of polymer segments on these sites releases the entropic forces of depletion and causes delamination of the polymer material from the surface. 7. CONCLUSION

Depletion from an impenetrable surface is a universal property of polymeric substances, which is potentially detrimental for the strength of adhesion bonds. Depletion is suppressed when the adsorption energy of polar groups in the polymer is large enough to compensate for the free energy of depletion. Adsorption reduces the effective distance H between the substrate and the polymer from the length A of a Kuhn segment to the size of an adsorbed monomer δ. The change ∆γ is the effect of adsorption too, but it is too small to explain why the presence of polar groups in the interface is so important for the strength of an adhesion bond. The reason why adsorption does increase the strength by more than an order of magnitude is that it suppresses the depletion effect and creates a closer contact of the polymer molecules with the substrate surface.


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REFERENCES 1. W.A. Zisman, in Contact Angle, Wettability and Adhesion, Adv. Chem. Series No. 43, p.1, Amer. Chem. Soc., Washington DC (1964). 2. F.M. Fowkes, in Contact Angle, Wettability and Adhesion, Adv. Chem. Series No. 43, p.99, Amer. Chem. Soc., Washington DC (1964). 3. G. Frens and A. van der Put, J. Adhesion. Sci. Technol. 12, 1355 (1998). 4. K.L. Mittal (Ed.), Physicochemical Aspects of Polymer Surfaces, Vols. 1&2, Plenum Press, New York (1982). 5. H.C. Hamaker, Physica, 4, 1058 (1937). 6. L.J. Evers, S.Yu. Shulepov and G. Frens, Faraday Discussions, 104 (1997). 7. L.J. Evers, S.Yu. Shulepov and G. Frens, Phys. Rev. Lett., 79, 4850 (1997). 8. J.F.M. Pennings, Colloid. Polym. Sci. 256, 1155 (1978). 9. J.F.M. Pennings and B. Bosman, Colloid. Polym. Sci. 257, 720 (1979). 10. J.F.M. Pennings, Colloid. Polym. Sci. 258, 1099 (1980). 11. J.F.M. Pennings, “Adsorption of Vinyl Copolymers from Melts and Solutions”, Ph.D. Thesis, TU Eindhoven (1982). 12. H.P. Poulis, in: “UV/Ozone Surface Treatment of Polymers for High-Quality Adhesion in the Footwear and Coating Industry”, Synthesis report (for publication), Brite Euram III-project no. BE-S2-2032, p. 9 (1998).



Attaining adhesion/cohesion within painted plastics ROSE A. RYNTZ∗ Visteon Corporation, 401 Southfield Road, Dearborn, MI 48121

Abstract—Adhesion to thermoplastic olefin (TPO) substrates, blends of polypropylene (PP) and an elastomer, is often difficult to attain due to the low surface energy, nonpolar nature of the surface. One often specifies, therefore, that the surface of the substrate be pretreated. Pretreatment methods can vary from oxidation of the surface, e.g., plasma or flame, to diffusion and mechanical interlocking of an applied polymer with the surface. Mechanical interlocking with the nonpolar TPO surface can be achieved through the use of an adhesion promoter, namely a chlorinated polyolefin (CPO). The type of CPO utilized, in addition to the types of solvents and heat effects utilized, can substantially influence the degree of adhesion/cohesion obtained within the CPO/TPO system. This paper reviews the factors influencing the adhesion/cohesion of painted TPO substrates. Heat histories, TPO molding variations, CPO types, including solvent and resin variations, and topcoat (basecoat/clearcoat) chemistries are all found to influence the adhesion/cohesion of the painted TPO assembly. Keywords: Adhesion; cohesion; plastic; chlorinated polyolefin; thermoplastic olefin.

1. INTRODUCTION AND DISCUSSION

The lack of adhesion to low surface energy substrates (surface energy ~ 30 mJ/m2 for PP) has been attributed to their inertness, including poor wettability, lack of polar functionality, and good solvent resistance. The introduction of an elastomer, producing a TPO, generally provides better adhesion than PP alone. Attaining adhesion to TPO, however, is still often problematic. The adhesion of organic coatings to a TPO is strongly dependent upon the type of surface treatment the thermoplastic material receives prior to painting. One such treatment involves the application of a thin layer (7 to 14 µm) of an adhesion promoting primer that is comprised of a CPO. Although materials of this type are known to establish good adhesion between the TPO and subsequent paint layers, the mechanism of adhesion promotion is not unequivocally known. The mechanism by which CPO coatings promote adhesion to TPO has been speculative over the last several years [1-5]. Waddington and Briggs [5] have reported experiments that suggest CPO promotes adhesion by diffusing into the ∗

Phone: (313)755-6164, Fax: (313)755-0601, E-mail: [email protected]

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R.A. Ryntz

substrate and forming a chain entangled interphase between substrate and adhesion promoter. They used secondary ion mass spectrometry (SIMS) and x-ray photoelectron spectroscopy (XPS) to study the peel failure between a CPO and two grades of TPO: one that exhibited strong adhesion and one that exhibited weak adhesion. Chlorine was detected on the surface of the TPO that exhibited strong adhesion, but not on the surface of the TPO that exhibited weak adhesion. In a separate experiment, they coated the two grades of TPO with a CPO and measured the amount of chlorine at the air interface of the CPO. Lower levels of chlorine were found over the TPO grade that exhibited better adhesion. This depletion of chlorine supported a mechanism based on the diffusion of CPO into the TPO. Clemens et al. [1] utilized surface characterization techniques, including electron spectroscopy for chemical analysis (ESCA), energy dispersive X-ray analysis (EDXA), time-of-flight secondary mass spectrometry (TOFSIMS), and transmission electron microscopy (TEM), in attempts to elucidate the mechanism by which CPO primer coatings promoted adhesion of paints to polypropylene and TPO. The coatings, their interphasial failures, and taper-cross sections were studied, using both waterborne and solventborne CPO primers. It was determined that CPO primers did not penetrate deeply into the polyolefin substrates, but were quite mobile following application of the topcoat. The mechanism by which CPO primer coatings promoted adhesion was discussed in terms of better surface wetting and thus contact, allowing for sufficient dispersion interactions with the substrate, and intermixing/bonding with the more polar topcoat. Solventborne CPOs generally showed failure at the CPO/polyolefin interface when dried at ambient temperatures. Waterborne CPOs showed equivalent adhesion promotion, although cohesive failure within the primer layer was more common and appeared to be a function of formulation parameters. Work done by Ryntz and coworkers [6, 7] demonstrated that adhesion to TPO was strongly influenced by the type and amount of solvent contained within the adhesion promoter (CPO containing primer) as well as the surface morphology of the TPO substrate. The morphological changes at and near the surface of TPO affect not only paint adhesion but also the cohesive integrity of the painted plastic assembly. Morphological changes in the TPO substrate are accomplished in the presence of solvent from the adhesion promoter (as well as from subsequent topcoats) and vary depending upon bake time and temperature. If it is assumed that the morphology of TPO is indicated by a closely packed, crystalline surface layer of PP, under which lies a rubber-rich boundary layer, then adhesion attainment is dependent upon the diffusion of CPO adhesion promoter through the crystalline PP surface layer and its swelling and entanglement with the rubber-rich layer beneath it (Figure 1). In general, a greater degree of rubber swelling is indicative of a higher amount of solvent penetration. Initial adhesion to TPO appears to be a function of polypropylene crystallinity at the surface only as it relates to the ability of solvents to penetrate through and into the rubber-rich layer beneath it.

Attaining adhesion/cohesion within painted plastics

31

Figure 1. Painting TPO.

One must be aware, however, of the rubber swelling effect on the ability of the TPO to withstand stress (compressive or shear) within its boundary layers in a swollen state. For example, when the rubber is swollen (elongated) to near its tensile break point, it would require little imparted stress to delaminate the rubbery boundary layer from the bulk morphology of the TPO beneath it. Peel strengths (given in terms of cohesive integrity values) performed on TPO plaques exposed to a compatible solvent (xylene, solubility parameter near that of TPO) indicated that those plaques exposed to solvent (30 minutes at 25°C) afforded lower values than those not exposed to solvent (Figure 2) [7]. This variance can be accounted for by increased swelling or plasticization of the top TPO layers. When the same plaques were swollen by slow evaporating (Aromatic 100) versus fast evaporating (tolu-

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Figure 2. TPO cohesive integrity as a function of solvent exposure.

ene) solvents and then flashed-off through the use of an infrared oven, very little difference in peel strength values was seen. In order to attain maximum adhesion to TPO, therefore, one tries to maximize solvent diffusion into the substrate while minimizing the amount of solvent remaining in the substrate after bake. Solvent diffusion, however, is only indicative of the ability of the CPO, the actual adhesion promoting moiety, to be carried into the substrate and not of the actual CPO penetration. In order to understand the amount of diffusion accomplished by both the CPO and solvent a fluorescent “tagging” experiment was performed to monitor the depth of interpenetration of both substances via fluorescence microscopy. The use of an unbound fluorescent dye, 7,7'-dimethyl-4aminocoumarin, in the solvent was utilized to follow the depth of solvent penetration. A covalently bound dye, attached directly to the maleic anhydride component of the CPO, was utilized to follow the depth of CPO penetration. This technique is not unfounded, being studied by Li et al. [8] to determine the surface morphology of a polymer blend through laser confocal fluorescence microscopy. The ability of the solvent to swell the TPO substrate (Figure 3) was studied in relation to the depth of solvent penetration into the substrate versus the bake temperature (Figures 4 and 5). This was accomplished through the use of fluorescence microscopy as described above. The depth of penetration of the “free” dye, 7,7'-dimethyl-4-aminocoumarin, was measured in a 70/30 PP/Rubber w/w compounded TPO as a function of bake temperature while maintaining equivalent solvent composition (80/20 xylene/Aromatic 100 based on weight). The dye penetration, as viewed from a cryogenically “cracked” cross section of the treated TPO substrate, was found to increase as a function of bake temperature (Figure 5 vs. Figure 4). The depth of penetration was followed by monitoring the


Figure 3. Swellability of TPO.

Figure 4. Free dye penetration into TPO at 25°C.

Figure 5. Free dye penetration into TPO at 121°C.

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Figure 6. Topographical fluorescence of “tagged CPO” on TPO.

increase in the optical intensity of the dye fluorescence (right hand side of both Figures). Dye penetration values were taken at 50% of the intial intensity of the fluorescence trace shown in the right side of the Figures. At room temperature, the dye is shown to migrate roughly 150 µm into the substrate, whereas at 121°C, the depth of penetration increases to roughly 600 µm. This depth of penetration merely mimics the amount of free dye carried into the substrate by the solvent and not the actual diffusion of the much higher molecular weight CPO molecule. In a separate fluorescence experiment the fluorescent tag was covalently bound directly to the maleic anhydride functionality of the CPO molecule. The penetration of the CPO molecule was then determined by monitoring the fluorescence with a laser confocal microscope. In Figure 6 one can view the CPO/TPO topographical fluorescence as a function of temperature. A surface profile of the orientation of fluorescence in the same section of the substrate as a function of bake temperature is displayed. It can be seen that as the fluorescently tagged CPO adhesion promoter is baked the fluorescent (light sections of the diagrams) regions appear to coalesce. One can postulate that the CPO diffuses through the TPO surface and entwines with the rubberrich areas of the TPO directly beneath the surface. As the swelling continues, through the interaction of the solvent-laden CPO with the rubber, the CPO/rubber regions begin to break through the surface and coalesce.


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Figure 7. Chloride ion diffusion into TPO via TOFSIMS analysis.

Laser confocal fluorescence microscopy also confirmed depths of penetration of the fluorescently tagged CPO. The depth of penetration was determined by rastering the laser across the selected section of the TPO at various focal planes. The depth of CPO penetration was then re-created by addition of the various focal plane depths where fluorescence was seen. The fluorescence penetration after application of the fluorescently tagged CPO and a 30-minute dwell time at 25°C was approximately 4 µm. After a 121°C bake, the depth of fluorescence increased to ca. 10 µm. Time-of-flight secondary-ion-mass spectrometry was utilized to evaluate the depth of penetration of the adhesion promoter based on imaging of the chloride ion contained within the CPO. The chloride images and line scans of crosssectioned samples of the chlorinated polyolefin based adhesion promoter coated TPO system are shown in Figure 7. The pixel intensity of the chloride ion (as determined with the gallium laser) can be viewed pictorially on the left side of the Figure while the absolute intensity is shown graphically on the right side of the Figure. A number of important features are immediately apparent in the images. A well-defined line, representing the adhesion promoter layer, is readily apparent in the pictorial side of the Figure. The thickness of this layer is about 10-15 µm. Localized areas of high intensity chloride observed along this sharp line are suggestive of topographical anomalies introduced into the sample surface by the microtoming technique used to prepare the sample. This turns out to be a critical factor in determining the degree of useful information that can be gained from the analysis. Most importantly, chloride diffusion into the TPO is noted graphically on the right side of the Figure. The diffusion into the TPO substrate appears to be on the order of 3 to 9 µm. The diffusion of a chlorinated polyolefin polymer having a molecular weight of about 15,000 could reasonably be expected to result in several orders of magnitude less diffusion into the TPO substrate. However, non-Fickian behavior resulting from solvent induced swelling, crazing and porosity induced by phase separation of the elastomer from the polypropylene would not be unexpected.

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Figure 8. Effect of CPO resin type on its adhesion (peel strength) to TPO.

These results correlate quite well with those determined by Komatsu et al. [9] as well as the TOF-SIMS studies. The slightly higher degree of interpenetration of the CPO into TPO, as suggested by the fluorescent studies relative to the SIMS data, can be explained in the context of molecular weight. A gel permeation chromatography (GPC) examination of the tagged CPO showed small amounts of low molecular weight moieties present within the tagged CPO. The low molecular weight components in the CPO compound are believed to contribute to the deeper depths of penetration. The effects of structural differences within the chlorinated polyolefins and the effect they have on coating performance is described in the work of Fujimoto [10] and Ryntz [11]. In effect, the work by the above-mentioned authors shows the effect of chlorination of the polyolefin on adhesion and crystallinity of the polymer. In summary, the higher the degree of chlorination the less crystalline is the chlorinated polyolefin, thus making it more compatible with other typical filmforming compounds used in coatings. However, the lower the level of chlorine the better the adhesion to olefinic substrates. These two opposing parameters must be balanced with one another to provide adhesion promoters with the best balance of properties. In the work performed by Ryntz [11], the degree of maleation of the chlorinated polyolefin, in conjunction with the chlorinated polyolefin molecular weight and molecular weight distribution, was also shown to affect the degree of adhesion between the chlorinated polyolefin and the TPO substrate (Figure 8). As can be seen in the figure, the resins with the highest 90 degree peel strength (triangles are indicative of relative peel strength only) were those with highest chlorine levels, some level of maleic anhydride, low number average molecular weight, low polydispersities, and low viscosities (CPO resins labeled “A” and “B”).


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Figure 9. Effect of molding condition on TPO swellability.

Although the work described above identified the importance of CPO structural attributes in adhesion of CPO to TPO, another anomaly was found to occur with respect to adhesion and location along the TPO substrate. As shown in Figure 9, when swelling experiments were performed near or away from the proximity to the “gate” (the location in a injection molded substrate where the hot plastic melt is introduced into the part), variations in swellability occurred [6]. As would be expected, those materials with more rubber should swell more. However, it was noted that the areas “near” the gate in the TPO part exhibited lower amounts of swelling. In a separate experiment, the adhesion of the topcoat “near” and “away” from the gate in TPO injection molded plaques was determined via a gas soak test [12]. It was found that when white topcoated, black pigmented TPO substrates were cross-hatched and immersed in ethanol laden gasoline, the adhesion began to become differentiated at the different locations. As shown in Figure 10, the white topcoated TPO substrates, when exhibiting adhesion loss, became black (indicative of the black substrate). In traversing from the top of the Figure to the bottom, the TPO flexural modulus increases and the fluorescent dye penetration decreases. With each set of TPO materials, there are portions of the substrate representing the “near” gate and “away” from the gate locations (there are four TPO materials, and each set of 2 represents the gate location). As is evidenced in the Figure, the lower modulus (higher rubber content) TPO materials exhibit better adhesion than those with higher modulus. As one traverses the Figure from top to bottom, it is also evident that the area “near” the gate in higher modulus substrates begins to fail before the area “away” from the gate. Several researchers have examined the importance of the variation in TPO surface and its affect on adhesion [7, 13, 14]. The surface morphology of semi-

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Figure 10. Gasoline resistance of various modulus topcoated TPO substrates as a function of position from gate.


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Figure 11. Birefringence vs. gate location in injection molded TPO.

crystalline polymers has been shown to be similar to or substantially different from that present in the bulk as a consequence of processing conditions [15]. The relatively minor changes in morphology lead to specific changes in the chemical, physical, and mechanical properties of the surface and of the interior or bulk. Since chemical reactivity, gas permeability, adhesion, friction, and abrasion resistance are dependent on surface morphology, it is of interest to delineate those factors that influence the formation of surface structure. It has been found that cross-polarized light can be utilized as an effective means with which to measure the skin/shear layer in injection molded TPO [7]. The shear zone, under cross-polarized light, is birefringent. As shown in Figure 11, however, the shear layer thickness can vary as a result of the way a part is filled. Near the injection molding gate, the shear layer appears thinner (57 µm) versus away from the gate (119 µm). It was also determined that the skin layer near the gate exhibited a greater “microhardness” than that away from the gate, possibly as a result of increased packing density or variations in surface crystallinities. The surface hardness variation may account for the adhesion variations, since it was also determined that areas “near” the gate exhibited lower dye penetration values. Fujiama and Kimura [13] found that a relationship existed between the thickness of the skin layer in polypropylene and its yield strength. The skin-layer is composed of a “shish-kebab” like structure where crystalline lamellae align perpendicular to the flow direction (referred to as transcrystallinity). Generally thicker shear zone skin layers afforded higher tensile strengths (Figure 12).

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Figure 12. Yield strength of TPO as a function of skin layer thickness [14].

The types of materials that contribute to the “skin layer” in injection molded TPO can vary [7]. If the TPO contains low molecular weight additives, e.g., processing stabilizers, viscosity adjusters, or even process degraded or low molecular weight “tails” of the polymers, one can experience additive migration to the TPO surface (Figure 13). The species migrating to the surface can amount to formation of a “weak boundary layer (WBL)” that exhibits poor cohesion to the bulk TPO. The relationship of “skin layer” thickness to cohesive integrity of the TPO substrate became an important consideration when a topcoated substrate was exposed to a compressive/shear stress [16]. Thermoplastic polyolefin substrates, under compressive tensile shear stress, often delaminate cohesively within the TPO substrate, where cohesive delamination was found to be dependent on load, temperature, and shear velocity parameters of the tensile shearing event. The failure was found to occur in the weak boundary layers of the TPO substrate, most often between 25 and 100 µm into the substrate (Figure 14). It is interesting to note that if one exposes unpainted TPO to the compressive/tensile shearing, no failure occurs. It appears that the coefficient of friction of the surface, being very low for polypropylene, allows the shear load to be transmitted across the monomolecular low friction layer. When painted, however, the friction coefficient of the paint, being higher than that of the TPO surface, allows the force to be transmitted into the painted TPO surface.


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Figure 13. Formation of weak boundary layers within TPO as a function of heat.

Figure 14. Gouge in topcoated TPO.

It was determined that the most significant parameters in the paint which affected “gouge” damage resistance were glass transition temperature, secant modulus at break, and the static coefficient of friction [17]. In order to better understand, however, the attributes of the coating type utilized in painting TPO in relationship to the gouge resistance of the system a model coating system was de-

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Figure 15. Gouge performance vs. clearcoat type.

signed [16]. Waterbased hydroxyl functional styrenated acrylic clearcoats, containing non-aggressive solvents so as not to swell the TPO, were formulated. The gouge resistance of the formulated clearcoats was measured on a compressiveshear loading device. Figure 15 depicts the gouge resistance of the clearcoats. The gouge area/energy rating for the most severe test conditions, labeled DOE Run #8 (68.3°C, sliding velocity of 3.4 cm/sec, and a load of 272.4 kg) versus the less severe conditions (labeled DOE Runs #2, #4, and #6, where the loads, sliding velocities, and temperatures are lower) are shown. The smaller the value (a rating of 0 depicts no damage) the better the performance. The coating formulation nomenclature is: #77A, D, G, and J are hydroxyl functional styrenated acrylics with a Tg of 45°C; where 77D and J contain surface free energy modifier where 77 A and D are not crosslinked where 77 G and J are crosslinked with a blocked isocyanurate #77C, F, I, and L are hydroxyl functional styrenated acrylics with a Tg of 65°C; where 77 F and L contain surface free energy modifier where 77 C and F are not crosslinked where 77 I and L are crosslinked with a blocked isocyanurate As viewed in Figure 15, all of the coatings, regardless of the Tg or crosslinker level, when formulated with added surface free energy modifier (mar agent), showed no gouge (77-D, 77-J, 77-F, and 77-L). It appears generally that as the Tg increased (77-C vs. 77-A and 77-I vs. 77-G) the gouge damage increased. It is be-


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lieved that the brittleness of the higher Tg resin accounts for brittle fracture of the coating during applied compressive load resulting in asperity formation and subsequent gouging. Added crosslinker, 77-G vs. 77-A and 77-I vs. 77-C, appeared to improve gouge resistance. This is believed to be due to an increase in tensile strength of the coating when crosslinked. 2. SUMMARY

Mechanical interlocking of topcoat with the nonpolar TPO surface can be achieved through the use of an adhesion promoter, namely a chlorinated polyolefin (CPO). The type of CPO utilized, in addition to the types of solvents and heat effects utilized, can substantially influence the degree of adhesion/cohesion obtained within the CPO/TPO system. Heat histories, TPO molding variations, CPO types, including solvent and resin variations, and topcoat (basecoat/clearcoat) chemistries were all found to influence the adhesion/cohesion of the painted TPO assembly. REFERENCES 1. R.J. Clemens, G.N. Batts, J.E. Lawniczak, K.P. Middleton and C. Saas, Proc. XIXth Conf. on Organic Coatings Science and Technology, Athens, Greece, p. 105 (1993). 2. K. Grundke and H.J. Jacobasch, Farbe und Lacke, 98, 934 (1992). 3. Y. Aoki, J. Polym. Sci.: Part C, 23, 855 (1968). 4. T.J. Prater, S.L. Kaberline, J.W. Holubka and R.A. Ryntz, J. Coatings Technol., 68 (857), 83 (1996). 5. S. Waddington and D. Briggs, Polymer Commun., 32 (16), 506 (1991). 6. R.A. Ryntz, Q. Xie and A.C. Ramamurthy, J. Coatings Technol., 67 (843), 45 (1995). 7. R.A. Ryntz, Prog. Org. Coatings, 27, 241 (1996). 8. L. Li, S. Sosnowski, C. Chaffey, S. Balke and M. Winnik, Langmuir, 10, 2495 (1994). 9. Y. Komatsu, J. Kamimura, O. Aoki, D. Chung and M. Wiseman, SAE Technical Paper #940187 (1994). 10. F. Fujimoto, Paint and Resin, 36 (February 1986). 11. R.A. Ryntz, Advanced Coating Technology Conference Proceedings, Engineering Society of Detroit, Detroit, MI (April 1997). 12. R.A. Ryntz and B. Buzdon, Prog. Organic Coatings, 32, 167 (1997). 13. M. Fujiyama and S. Kimura, Kobunshi Kobunshu (English Edition), 4 (10), 777 (1975). 14. M.R. Kantz, H.D. Newman and F.H. Stigale, J. App. Polym. Sci., 16, 1249 (1972). 15. D.R. Fitchmun and S.J. Newman, J. Polym. Sci.: Part A2, 8, 1545 (1970). 16. R.A. Ryntz, M. Everson and G. Pollano, Prog. Organic Coatings, 31, 281 (1997). 17. B. Buzdon and R.A. Ryntz, SAE Paper, presented at SAE Conference, Detroit, MI (1996).



Scanning electric potential microscopy (SEPM) and electric force microscopy (EFM) imaging of polymer surfaces ELIZABETH F. DE SOUZA,1 MÁRCIA MARIA RIPPEL,2 AMAURI JOSÉ KESLAREK,2 ANDRÉ GALEMBECK,2 CARLOS ALBERTO RODRIGUES COSTA,2 ÉRICO TEIXEIRA NETO2 and FERNANDO GALEMBECK2,∗ 1

Instituto de Ciências Biológicas e Química, Pontifícia Universidade Católica de Campinas PUC-Campinas, PO Box 1111, CEP 13020-904, Campinas - SP, Brazil 2 Instituto de Química, Universidade Estadual de Campinas - UNICAMP, PO Box 6154, CEP 13083-970, Campinas - SP, Brazil

Abstract—Scanning electric potential microscopy (SEPM) and electric force microscopy (EFM) images of polymer surfaces are presented and compared to standard non-contact AFM images, acquired concurrently for the same areas. Samples used were synthetic latex films with well-known distribution of chemical constituents and thus of ionic electrical charges, as well as a natural latex film and two finished industrial samples. Pairs of AFM and SEPM or EFM images show a variable degree of correlation, thus evidencing the independence of topographical and electrical features of the samples, on the micro- and nanoscopic scales. In every case, contrasting domains with non-zero electric potentials are observed, extending for a few tenths of a micrometer and creating an electric mosaic in the otherwise neutral polymers. Large electric potential gradients across the film surfaces are observed which are probably relevant for polymer surface adhesion, coating and contamination. Keywords: Electric charges in polymers; scanning electric potential microscopy (SEPM); electric force microscopy (EFM); electric maps.

1. INTRODUCTION

Polymer surface properties are closely related to the performance of these materials in areas of adhesion and coating, biocompatibility, and contamination [1]. Polyolefins and related thermoplastics are often used as dielectrics and they have a strong ability to acquire electrical static charges especially at their surfaces. Many other important problems are related to the presence of electric charges in

∗

To whom all correspondence should be addressed. Phone: 055 51 788 3080, Fax: 055 51 788 3023, E-mail: [email protected]

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dielectrics, such as: electret formation, space and residual charges, double-layer formation at interfaces, and interfacial polarization [2]. There is a relevant conceptual constraint in dealing with these problems, which is the idea of charge neutrality as a normal state for polymer dielectrics at every relevant length scale, from the macromolecules up to the macroscopic plastic parts (films, coatings, tubes) and devices [3]. A direct measurement of electric potentials in dielectrics is a difficult task, although suitable techniques have been available for the past century. Specific techniques have been devised for the case of monolayers formed by organic substances (e.g. surfactants, polymers) on top of aqueous solutions [4]. For instance, the oscillating electrode technique has revealed a great deal of information on the distribution and orientation of different substances on liquid surfaces [5]. More recently, the advent of the scanning probe microscopes has made available techniques for producing maps showing local changes of charge density on insulator surfaces [6]. Other properties measured with spatial resolution by using these instruments are: dielectric constant [7], film thickness of insulating layers [8], surface potential [9] and ferroelectric susceptibility [10]. A scanning force microscope measures forces between the probing tip and the sample surface [11]. It is comprised of a force sensor mounted on a cantilever (tip), an optical sensor, a positioning feedback system, an x-y-z piezoelectric scanning system, and a microcomputer for the acquisition of data and image processing [12]. A scheme of the basic setup of an atomic force microscope (AFM) is presented in Figure 1. Many modes of operation for AFM have been developed, and non-contact mode AFM is the most common. In the dynamic non-contact operation, the cantilever is mechanically oscillated to vibrate on or near its resonance frequency by a piezoelectric vibrator. The gradient of the force between the tip and sample modifies the compliance of the cantilever, hence inducing a change in vibration amplitude as a function of the tip-sample spacing, which is in the 10-100 nm range. The associated circuitry monitors the resonant frequency or vibrational amplitude of the cantilever and keeps it constant, with the aid of a feedback system that moves the scanner up and down. By keeping the resonant frequency or amplitude constant, the system allows the tip-to-sample spacing to be held constant for profiling applications [12, 13]. Several forces contribute to the deflection of an AFM cantilever, but long-range van der Waals forces predominate in the non-contact mode [14]. Electrostatic forces can also deflect the cantilever when it scans over locally charged domains of the sample surface. This is caused by the induction of charge in the tip, which results in changes in the tip-to-surface force gradient [15]. For instance, over a negatively charged region of the sample surface, positive charges are induced in the tip, resulting in an additional attractive electric force. The same attraction occurs between a positively charged region of the sample and the negative charges induced in the tip. In these two cases, the force between the tip and

SEPM and EFM imaging of polymer surfaces

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Figure 1. Schematic diagram of the setup of an atomic force microscope operating in the noncontact mode. The inset shows a magnified representation of the tip-sample contact region.

the sample is increased and it is impossible to identify the electric charge signal in the area under examination. However, if a positive potential is applied to the sample-holder, the negative charges of the sample are compensated, resulting in a decrease of the attractive electric force between the sample and the tip due to a decrease in the induced charge. In the neutral regions of the sample, the positive potential applied to the sample holder induces negative charges in the tip, which is attracted to the surface. Finally, in the positive domains of the sample the tip is under the added electrostatic attraction of the sample and the sample holder. Therefore, the use of an external potential applied to the sample-holder allows us to identify the charge signal in the sample surface. Figure 2 presents a schematic representation of charges induced on the scanning tip, and the corresponding profiles of the attractive electrostatic force acting on the tip. In this work, we used non-contact atomic force microscopy (AFM) to identify the morphological features of various polymeric surfaces and electric force microscopy (EFM) as well as scanning electric potential microscopy (SEPM) to investigate the microelectric characteristics of the examined surfaces. These two techniques for electric sensing can be practiced using available scanning force microscopes, but their use is still very limited, as can be gathered from the literature and relevant databases.

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Figure 2. Schematic representation of charges induced on the scanning tip, and the corresponding profiles of the attractive electrostatic force acting on the tip. (a) The sample holder is electrically neutral: the electric force between the tip with induced charges and the sample surface charged domains are always attractive. (b) The sample holder is positively charged by the application of an external electric potential: the attractive electric force between the tip and the sample surface is reduced over the negative domains of the sample surface, which are compensated by the positive charges of the sample holder. (c) The sample holder is negatively charged by the application of an external electric potential. (d) The negative charge of the sample holder is increased by the application of a more negative external electric potential.

2. EXPERIMENTAL

2.1. Preparation of samples Titanium dioxide: This sample was used as a control, to verify the separation of van der Waals and electrostatic contributions to imaging. A drop of a very dilute aqueous dispersion of titanium dioxide nanometric particles (Riedel-de Haën AG, Seelze-Hannover, Germany) was placed on freshly cleaved mica and allowed to dry in air. Poly(styrene-co-2-hydroxyethylmethacrylate) (PS-HEMA) latex: The PSHEMA latex was prepared by batch surfactant-free emulsion copolymerization of styrene and 2-hydroxyethylmethacrylate. Details on the latex preparation and characterization are given in previous works from this laboratory [16, 17]. The


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PS-HEMA latex film was prepared by depositing of one droplet of the latex dispersion (1% solids) directly on top of a freshly cleaved mica surface, and allowing to dry at room temperature for 48 h. Poly(styrene-co-acrylamide) (PS-AAM) latex: The PS-AAM latex was prepared in this laboratory by batch surfactant-free emulsion copolymerization of styrene and acrylamide [18]. The PS-AAM latex was dialyzed and further fractionated by zonal centrifugation in a linear aqueous sucrose density gradient (100µL latex, 1% solids, 19 krpm). Each latex fraction was then dialyzed to remove sucrose by centrifugal ultrafiltration [19], and the denser latex particles were deposited on mica, as in the case of PS-HEMA. Natural latex (NL): Raw latex from Hevea brasiliensis (from a Rubber Research Institute of Malaysia 600 clone) without any stabilizing agent was collected at the Instituto Agronômico de Campinas and centrifuged at 10 krpm, for 2 h at 5°C. One droplet of the latex dispersion (2.6% solids) was dried on top of a freshly cleaved mica surface at room temperature for 48 h, thus yielding a film. Poly(styrene-co-n-butylacrylate) (PS-BA) latex: Droplets of the non-dialyzed latex dispersion (1% solids) were deposited directly on top of a glass slide cover and dried at (20 ± 2)°C, under 50% relative humidity, for 48 h. Commercial RetroplanTM polyester film (Polyester): Small fractions of the printable and of the back sides of a polyester film were fixed with carbon-filled adhesive onto a freshly cleaved mica surface. Commercial high-density polyethylene film (HDPE): The HDPE film was washed with water, ethanol and air-dried. Then a small fraction of the film was also fixed with a conductive adhesive onto a freshly cleaved mica surface. 2.2. Atomic force microscopy (AFM) The dry latex and commercial films obtained as described above were examined by AFM in order to observe their morphological features. AFM measurements were performed in a Topometrix DiscovererTM instrument operating in noncontact mode, under air and concurrently with EFM or SEPM imaging. Topography changes were sensed by monitoring the detector signal amplitude at 300 x 300-pixel resolution, using platinum-coated silicon nitride tips with 20 nm nominal radius. This is one-third or less of the pixel size, in the images presented in this paper, and we avoided considering image features smaller than 60 nm. The scan rate was 2.5 µm/s. 2.3. Electric force microscopy (EFM) To perform EFM measurements, each sample line was scanned at two heights above the surface: 10 nm and then 60 nm. In the first scan (10 nm) the tip response is dominated by short-distance van der Waals forces. In the second scan (60 nm) the electrostatic interactions predominate, due to their slower dependence on distance. As the tip crosses over an electrically charged region, charge is induced in the tip resulting in changes in the tip-to-surface force gradient, causing a

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change in the effective cantilever spring constant and in the resonance frequency of the tip. This change in the resonance frequency modifies the signal phase detected by a lock-in amplifier. Therefore, the EFM signal is measured by monitoring the changes in the signal phase [20]. Figure 3 shows a scheme of the basic setup of an atomic force microscope used to perform electric force microscopic measurements. Both EFM and AFM measurements were made simultaneously, by monitoring the detector signal amplitude at 300 x 300-pixel resolution, using silicon nitride

Figure 3. (a) Schematic diagram of the TopometrixTM setup for simultaneous non-contact topographic and EFM measurements. (b) The force versus distance plot indicates the different forces acting on the tip and their relative magnitudes at different tip-sample distances. Note that when the distance between the tip and the sample surface is higher than d0, the electrostatic attractive force predominates. (c) The schematic diagram shows the tip-sample distances to obtain topographic (d < d0) or electric force data (d > d0).


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tips coated with platinum with 20 nm nominal radius (resonance frequency = 7099 kHz, stiffness constant = 1.8-5.2 N/m). 2.4. Scanning electric potential microscopy (SEPM) A common method to measure the contact potential difference between two materials is the vibrating capacitor or Kelvin method. The principle of the SEPM is similar to the Kelvin method except that forces are measured instead of currents [21, 22]. Scanning electric potential microscopy (SEPM) is then performed by applying a DC voltage between the tip and the sample while the cantilever hovers above the surface, without touching it. The cantilever deflects when it scans over the locally charged domains of the sample surface, and the magnitude of the deflection detected by the displacement of the sensor is proportional to the charge density [14, 15, 23]. In the case of SEPM, contrast is obtained when there are local electric potential gradients, across the sample surface. The images are obtained using an experimental arrangement in which the tip of the scanning force microscope is electrically biased as if it were a classical vibrating electrode in a monolayer surface potential measuring instrument. During the measurement, the resonant frequency of the cantilever is tracked by the four-quadrant photodetector and analyzed by two feedback loops. The first loop is used in the conventional way to move the scanner up and down, keeping a constant distance between the tip and the sample surface. The second loop is used to minimize the electric field between the tip and the sample. The second lock-in amplifier measures the tip vibration at the AC frequency oscillation while scanning, and adds a DC bias to the tip, to recover the undisturbed AC oscillation. The image is built using the DC voltage fed to the tip at every pixel [22]. The scheme of the basic setup of an atomic force microscope used to perform SEPM measurements is shown in Figure 4. SEPM imaging was done using the same tips and cantilevers as in AFM/EFM, and a DC voltage in the ±13 V range was applied between the sample and the tip of the microscope. 3. RESULTS AND DISCUSSION

The topographic image of titanium dioxide nanometric particles on mica (Figure 5 - top) shows that some of the particles are deposited adjacent to a step in the mica surface, the height of which is ca. 0.8 µm. In the SEPM image (Figure 5 - bottom) it is not possible to observe any difference in the electric force due to the mica step, but the sides of the oxide particles are perfectly visible. The outer layers of the oxide particle are positive, relative to the particle centers. This result shows an excellent separation of the contributions of the van der Waals forces and attractive electric forces acting on the tip. The electric map is thus completely independent of sample roughness.

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Figure 4. Schematic diagram of the TopometrixTM setup for SEPM measurements.

The PS-HEMA latex forms ordered arrays in both the liquid and the dry states, very easily [24]. Figure 6 shows the AFM and EFM images of a self-assembled macrocrystalline film formed by PS-HEMA latex. In the EFM image, the particle shells appear positive, relative to their cores. While considering the EFM and SEPM images, we should recall that these are not strictly surface imaging techniques, because the electrostatic interactions are long-range. Consequently, the sampled depth is estimated as some tens of nanometers. The elemental distribution maps of these particles were previously obtained using energy-loss spectroscopy imaging in the transmission electron microscope (ELSI-TEM) [16, 17]. The dry particles contain sulfur-rich (from sulfate) cores and potassium-rich shells. Consequently, the EFM and ELSI-TEM results are in excellent agreement concerning the electrical characteristics of particle shells and cores. The presence of ionic charges in a latex is well acknowledged, but their location in the dry films has not been examined in detail, up to now [25]. The present results show that ions are clustered in dry latex, forming domains with excess charge, extending for tens of nanometers.


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Figure 5. AFM (top) and SEPM (bottom) images of titanium dioxide nanometric particles on mica. The AFM image was rotated 90° counterclockwise and redrawn in perspective, to facilitate the observation of the TiO2 particles adjacent to the mica step.

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Figure 6. AFM (top) and EFM (bottom) images of the self-assembled macrocrystalline film formed by the poly(styrene-co-2-hydroxyethylmethacrylate) (PS-HEMA) latex.


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The AFM and SEPM images of the self-assembled film formed by poly(styrene-co-acrylamide) (PS-AAM) latex are presented in Figure 7. Particle cores are negative in relation to particle shells, as in the previous case of PSHEMA. This agrees with the known elemental distribution in these particles, as determined by ELSI-TEM [18]. In the natural latex film the situation is completely different: the topography image (Figure 8 - top) shows a continuous smooth film (maximum height = 12 nm) and without marked morphological features that could be identified as the original latex particles. On the other hand, the SEPM image (Figure 8 - bottom) shows a large contrast between neighboring domains, which can be associated with individual (darker, negative) latex particle cores, immersed in an electrically positive (bright) continuum. The SEPM image also shows some small brighter spots within the darker domains, which are probably related to latex non-rubber components [26]. Figures 9 and 10 present AFM and SEPM images of the surfaces of two regions of a film formed by a low Tg poly(styrene-co-n-butylacrylate) (PS-BA) latex. The topography images (Figure 9 and 10 - top) show continuous smooth films (maximum height = 34 nm), but with distinct morphological features. In both areas examined, there is correlation between the SEPM image and the topography observed, but in Fig. 9 it is possible to observe a sharp electric contrast between neighboring areas of the same morphological domain while the background is almost homogeneous. On the other hand, the topographic image in Fig. 10 presents roughly elliptical, negative domains, and an undulated background. Figure 11 presents the shadowed-AFM and EFM images of a high-density polyethylene film. The EFM contrast is very high, and both the upper and lower limits of the bias voltage scale were reached in a large number of pixels, which means that some parts of the image were saturated. Generally speaking, the surface protuberances appear as more negative than the depressions, but there are some areas in which this type of correlation is not observed. This is in agreement with observations for other polyolefins [27]. Both sides of a commercial polyester desk-jet printing film were imaged, as presented in Figures 12 and 13. At the printing side of this film there are many fuse-shaped elevations, pointing towards the right side of the picture. In the EFM image (Figure 12 - bottom), these elevations appear as if they were made of two parts, one of which to the left is more negative than the other. The printing side of this film is coated with a hydrophilic material, and the observed elevations are thus part of this material. The shape of the elevations suggests that the coating droplets spread during application, and the observed separation into domains of different electrical nature may be understood assuming that the coating is, e.g., made of a polyelectrolyte solution, in which the solvent liquid (or the serum) has a charge opposite to the polyelectrolyte. The backside of the polyester film is very smooth, both topographically (Figure 13 - top) and electrically (Figure 13 - bottom). Tiny shallow depressions appear as positive domains, but there is no correlation between the background topography and electrical contrast features.

56


Figure 7. AFM (top) and SEPM (bottom) images of the self-assembled film formed by poly(styrene-co-acrylamide) (PS-AAM) latex.


57

Figure 8. AFM (top) and SEPM (bottom) images of the film formed by the natural latex from Hevea brasiliensis (collected from a Rubber Research Institute of Malaysia 600 clone, film thickness = 250 nm).

58


Figure 9. AFM (top) and SEPM (bottom) images of the film formed by the non-dialyzed poly(styrene-co-n-butylacrylate) latex (film thickness = 185 µm).


59

Figure 10. AFM (top) and SEPM (bottom) images of the film formed by the non-dialyzed poly(styrene-co-n-butylacrylate) latex (film thickness = 185 µm).

60


Figure 11. AFM (top) and EFM (bottom) images of a high-density polyethylene film (film thickness = 93 µm). The AFM image was pseudo-illuminated at 90°, to facilitate the observation of the sample relief.


61

Figure 12. AFM (top) and EFM (bottom) images of the printable side of a commercial RetroplanTM polyester film (film thickness = 115 µm).

62


Figure 13. AFM (top) and EFM (bottom) images of the backside of a commercial RetroplanTM polyester film (film thickness = 115 µm).


63

4. CONCLUSIONS

The SEPM and EFM images of the various polymer samples examined reveal strongly contrasting image domains, with variable degrees of correlation with the respective topography images. This variability in the correlation between the imaging modes (topographic and electric) shows that there is a variable correlation between topographical and electrical features. The absence of correlation observed in some cases shows that these images are formed by independent factors. On the other hand, there is sufficient correlation in many image pairs, to allow for the identification of electrical features with corresponding topographical features. Generally speaking, the EFM images show more narrow features than SEPM, which may be an indication for the need to improve the feedback loop of the electric signal in SEPM. Of course, the detailed interpretation of all the systems examined in this work cannot be done without further additional characterization data. However, we believe that presenting these results at this stage may help motivate others to use the imaging capabilities of EFM and SEPM and other scanning probe techniques, and to develop the required interpretation tools. Progress is also required to obtain local electric charge densities from the electric force and electric potential maps, analogous to results already achieved in other force microscopy techniques, e.g. in topography-adhesion-stiffness imaging [28]. Finally, the present results show that neutral dielectrics are indeed made out of domains of different electrical nature, extending from a few nanometers to micrometers. Consequently, polymer electroneutrality is a result of charge balance but on a supramolecular or colloidal scale, not necessarily on the size scale of ion pairs or ion clusters. Acknowledgements The authors thank FAPESP, Pronex/Finep/MCT and CNPq. REFERENCES 1. S.-B. Lee, in: Polymer Interfaces and Emulsions, K. Esumi (Ed.), pp. 33–54. Marcel Dekker, New York (1999). 2. P. Robert, Electrical and Magnetic Properties of Materials. Artech House, Norwood, MA (1988). 3. K. Wu, M.J. Iedema and J.P. Cowin, Science, 286, 2482–2485 (1999). 4. G.L. Gaines, Jr., Insoluble Monolayers at Liquid-Gas Interfaces. Wiley-Interscience, New York (1966). 5. D.J. Shaw, Introduction to Colloid and Surface Chemistry. Butterworths, London (1966). 6. T. Hidaka, T. Maruyama, M. Saitoh, N. Mikoshiba, M. Shimizu, T. Shiosaki, L.A. Wills, R. Hiskes, S.A. Dicarolis and J. Amano, Appl. Phys. Lett. A 68, 2358–2359 (1996). 7. Y. Martin, D.W. Abraham and H.K. Wickramasinghe, Appl. Phys. Lett. 52, 1103–1105 (1988). 8. M. Nonnenmacher, O. Volter, J. Greschner and R. Kassing, J. Vac. Sci. Technol. B 9, 1358– 1362 (1991).

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9. J.W. Hong, S.-I. Park and Z.G. Khim, Rev. Sci. Instrum. 70, 1735–1739 (1999). 10. F. Saurenbach and B.D. Terris, Appl. Phys. Lett. 56, 1703–1705 (1990). 11. G. Binning, C.F. Quate and C. Gerber, Phys. Rev. Lett. 56, 930–933 (1986). 12. R. Lüthi, H. Haefke, K.-P. Meyer, E. Meyer, L. Howald and H.-J. Güntherodt, J. Appl. Phys. 74, 7461–7471 (1993). 13. Y. Martin, C.C. Williams and H.K. Wickramasinghe, J. Appl. Phys. 61, 4723–4729 (1987). 14. R. Howland and L. Benatar, A Practical Guide to Scanning Probe Microscopy, Park Scientific Instruments, Sunnyvale, CA (1996). 15. S. Belaidi, P. Girard and G. Leveque, J. Appl. Phys. 81, 1023–1030 (1997). 16. A.H. Cardoso, C.A.P. Leite, and F. Galembeck, Langmuir 14, 4447–4453 (1998). 17. A.H. Cardoso, C.A.P. Leite, M.E.D. Zaniquelli and F. Galembeck, Colloids Surfaces A 144, 207–217 (1998). 18. E. Teixeira-Neto, C.A.P. Leite, A.H. Cardoso, M.C.V.M. Silva, M. Braga and F. Galembeck, J. Colloid Interface Sci. 231, 182-189 (2000). 19. F. Galembeck and E.F. Souza, in: Polymer Interfaces and Emulsions, K. Esumi (Ed.), pp. 119– 164, Marcel Dekker, New York (1999). 20. H. Bluhm, A. Wadas and R. Wiesendanger, Phys. Rev. B 55, 4–7 (1997). 21. G. Ertl and J. Küppers, Low Energy Electrons and Surface Chemistry, Verlag Chemie, Weinheim (1974). 22. M. Nonnenmacher, M.P. O’Boyle and H.K. Wickramasinghe, Appl. Phys. Lett. 58, 2921–2923 (1991). 23. R.M. Nyffenegger, R.M. Penner and R. Schierle, Appl. Phys. Lett. 71, 1878–1880 (1997). 24. A.H. Cardoso, C.A.P. Leite and F. Galembeck, J. Brazilian Chem. Soc. 10, 497–504 (1999). 25. M. Braga, C.A.R. Costa, C.A.P. Leite, M.C.V.M. da Silva and F. Galembeck, J. Phys. Chem. B 105, 3005–3001 (2001). 26. M.R. Sethuraj and N.M. Mathew, Natural Rubber: Biology, Cultivation and Technology, Elsevier Science, Amsterdam (1992). 27. A. Galembeck, C.A.R. Costa, M.C.V.M. da Silva, E.F. de Souza and F. Galembeck, Polymer, 42, 4845–4851 (2001). 28. Y. Zhang, A. Pungor, G. Jogikalmath and V. Hlady, J. Adhesion Sci. Technol., 14, 1751–1765 (2000).


The residue (smut) formed on aluminum alloys during hydrofluoric acid etching and its effect on a coating process ANURAG P.S. TIHAIYA,1 JAMES P. BELL∗,1 and GUY D. DAVIS2 1

Institute of Materials Science, University of Connecticut, 97 N. Eagleville Road, Storrs, CT 06269-3136, USA 2 DACCO SCI Inc., 10260 Old Columbia Rd, Columbia, MD 21046, USA

Abstract—Hydrofluoric acid (HF) etching was studied as a pretreatment process for aluminum alloys. A coating process involving spontaneous polymerization on aluminum surfaces was used as a basis for the study. Five different aluminum alloys (2024, 3105, 5000, 6061 and 7075) were examined. The HF-treated aluminum surfaces were analyzed by means of XPS, AES and ESEM. HF etching resulted in differing behaviors among the alloys, in terms of extent of etching, surface morphology obtained after etching, and the quantity of residual black etching product on the surface. These factors led to different coating rates among the alloys. The residual etching product (“smut”) was observed to affect the polymerization rate on the surface and was necessary in some cases for coating formation. The smut consisted of hexagonal crystals and was identified as being primarily AlF3 · 3H2O. The presence of smut also affected the adhesion strength of the coating formed on the surface. Keywords: Hydrofluoric acid etching; aluminum pretreatment; residual product of etching; coating process.

1. INTRODUCTION

There is an abundant literature related to the surface pretreatment of aluminum. The subject has been the focus of numerous books and review articles [1-3]. However, studies of hydorofluoric (HF) etching of aluminum have not been as extensive. While some work has been done on the use of HF, alone or in solution with other acids, in aluminum pretreatment [4-7], the understanding of the effect of HF etching on various aluminum alloys is still limited. HF etching has been observed to result in formation of a residual product on the metal surface; some authors have reported formation of some kind of smut with alkaline [1, 8, 9] or acidic [2, 9] solutions. Little has been done to identify the smut material and un∗

To whom all correspondence should be addressed. Phone: (860) 486 4629, Fax: (860) 486 4745, E-mail: [email protected]

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A.P.S. Tihaiya et al.

derstand its role in subsequent processing. The present work involves a study of etching by HF of different aluminum alloys with a focus on the smut. A novel coating process (S-Poly©) involving Spontaneous Polymerization [10, 11] on treated aluminum surfaces was used as a basis for the study. 2. EXPERIMENTAL

2.1. Materials/chemicals Five different aluminum alloys (AAs) were used for this study. The alloys were 2024-T3, 3105, 5000, 6061-T6 and 7075-T6, supplied by United Technologies’ Sikorsky Aircraft division. AAs 2024 and 7075 were Al-clad. AAs 7075, 6061 and 2024 samples were 1.0 mm thick. The thicknesses of AA 3105 and 5000 were 0.6 mm and 0.3 mm, respectively. Aluminum coupons used in the coating process were 7.5 cm x 2.5 cm in size. They were cut from the source metal using a shear cutter in the machine shop at the University of Connecticut. BlueGold®, an industrial alkaline cleaner, manufactured by Modern Chemical Inc. was used for cleaning aluminum substrates. The cleaner contained diethylene glycol n-butyl ether, and had a pH of 13 in its concentrated state. A diluted solution of the cleaner in tap water (10% by vol.) was used for cleaning. The pH of the dilute solution was ~ 12. The aluminum samples were cleaned with a Kimwipes® tissue which had been dipped in the dilute solution of the cleaner, and then rinsed under flowing tap water. The reagent grade HF acid (48-50%) used for preparing the etch bath was purchased from Fisher Scientific Company. N-phenyl maleimide (NPMI) (TCI America), bismaleimide (BMI) (TCI America), styrene (Aldrich Chemical Company) and 2 (methacryloyloxy) ethyl acetoacetate (MEA) (Aldrich Chemical Company) were the chemicals used as monomers in the S-Poly© coatings. N-methyl pyrrolidinone (NMP) (Lancaster Synthesis, Inc., Pelham, NH, USA) and water (distilled using a Barnstead Model 210 Distilled Water Center) constituted the solvent system for the coating. The chemicals were used as received without any further purification. 2.2. Pretreatment The HF etch bath was prepared by diluting 48-50% HF with distilled water to the desired concentration. The concentration levels used were: 2.5%, 5.0% and 7.5% (which were obtained by diluting the original 48-50% HF by 20, 10 and 6.67 times, respectively). The volume of a typical etch bath (prepared and used in a polypropylene container) was 350-400 ml. The dip time in the etch bath varied from 20 seconds to 5 minutes. The following sequence of steps was followed for pretreatment of substrate coupons with HF solution. (1) Degreasing with a Kimwipes® tissue dipped in 10% solution of an industrial cleaner, BlueGold®.

The residue (smut) formed on aluminum alloys

67

(2) Washing thoroughly (by wiping with a wet tissue) under running cold tap water with moderate flow for 1-2 minutes. (3) Wiping dry with Kimwipes® tissue. (4) Immersing in the HF solution (using stainless steel tweezers). (5) Rinsing under running tap water with or without wiping the surface with wet Kimwipes® for 15-20 seconds. (6) Rinsing with distilled water (using a wash bottle). (7) Drying in an air flow, using rubber tubing connected to a filtered and compressed air supply, for 45 sec-1 min. The HF etching resulted in the formation of smut on the treated surface, with the amount varying as a function of alloy and the etching conditions. The smut was loosely adherent and could visually be removed from the surface by wet wiping (step 5). To study the effect of smut, some of the samples were left with the smut intact while the others were wiped; thus some of the samples were only rinsed after the etch, while the others were wiped with wet Kimwipes® tissue during rinsing. 2.3. Metal surface study Treated aluminum surfaces were analyzed by means of X-ray Photoelectron Spectroscopy (XPS), Environmental Scanning Electron Microscopy (ESEM) and Auger Electron Spectroscopy (AES). The XPS unit (Surface Science Instruments, Model SSX-100-03) used a monochromatic Al Kα source. The X-ray spot was 600 µm; an electron flood gun was used for charge neutralization. The survey spectra pass energy was 150 eV. The AES analysis (Physical Electronic Industries, Model PHI 25-110 spectrometer) was conducted using a LaB6 source. The ESEM unit featured a LaB6 source gun and a Gaseous Secondary Electron Detector (GSED). Some of the AA 6061 samples with smut were dried in a laboratory hood at ambient conditions after being removed from the etching bath. The smut was scraped off the surfaces with a thin polycarbonate blade and was analyzed by means of wide angle X-ray diffraction (XRD) and Inductively Coupled PlasmaAtomic Emission Spectroscopy (ICP-AES). The X-ray diffractometer (Bruker, D5005) used a copper tube (λ = 1.5418 Å). The smut powder was dissolved in an acidic solution (2 parts nitric acid (70%), 1 part hydrochloric acid (35%) and 1 part distilled water) at ambient conditions. The solution was further diluted with distilled water (1:75) for the ICP-AES analysis. 2.4. Studies on coatings The pretreated samples were coated by polymerization on the samples (coupons) surface via the S-Poly© method [10, 11]. The S-Poly© process is a novel coating method in which Spontaneous Polymerization forms a coating on a metal substrate immersed into a solution of monomers. The reaction initiates without the help of any external initiator, and the polymer is formed on the metal surface but

68


not in the solution. The metal surface is required to initiate the polymerization. The polymer forms on the metal surface with the polymer chains growing out from the surface instead of being pre-formed and deposited on it. This facilitates conformal and uniform coatings on substrates of various shapes, with excellent adhesion and minimal defects. The process has been studied in detail with respect to its application to aluminum [10], steel [11] and copper [12]. It is known to be able to apply to zinc, tin, silver and other metals as well. The pretreated samples were weighed and immediately placed in the polymerization bath. There was little time between the removal from the etch bath and coating of the samples (3-4 min). The samples were handled with stainless steel tweezers and placed on a Kimwipes® sheet for weighing. The conditions used for polymerization, extraction of residual monomers from wet coating (rinsing) and drying the coating were based on the results of prior work [10]. The volume ratio of NMP and water in NMP/water solvent system was kept at 57/43. Concentrations of the monomers in the S-Poly© solution were NPMI (0.2 M)/Styrene (0.2 M)/MEA (0.1)/BMI (0.005 M). The pH of the polymerization bath was adjusted to 3 by adding 0.5 M sulfuric acid solution. The polymerization was carried out at ambient conditions. Coated samples were rinsed in an aqueous solution of NMP (10% vol.) for 10 min to remove the residual monomers and NMP from the coating. Rinsed samples were placed in a laboratory oven with air circulation for drying. The oven was preheated to 110ºC and coatings were dried at this temperature for around 1-1.5 hour. Then the temperature was increased to 220ºC and maintained at this level for 0.75-1 hour. The samples were allowed to cool to room temperature in laboratory conditions. The coating rates on the samples were the basis for most of the comparative studies carried out in this work. The measured coating rate was actually based on a standard polymerization time of 10 minutes. Based on prior work on the kinetics of the polymerization [10], the coating rate for a shorter polymerization time (less than 10 min) is expected to be linear. At longer times the process becomes diffusion controlled and varies with the square root of time. First, the sample coupons were weighed before coating (after air drying) and, later, they were weighed after the coating was dried. The coated area was measured and the coating rate, which was based on the dry weight, was calculated in g/cm2/sec. The weight loss during etching was obtained by the difference of the weights before and after etching (air drying). Since the coating is formed as a result of the growth of polymer chains on the metal surface, the coating rate depends upon the number of the active sites (initiation sites) on the metal surface. This is discussed in detail in prior work [10]. The practical adhesion of the polymer coatings to the aluminum substrate was measured using the torsional testing method of Bell and Lin [13]. In the test, aluminum joints were used. The coating-metal bond was made to fail in shear and the bond shear strength was calculated. By visual examination, the failure occurred in the coating-metal interphase.


69

The substrates were etched and coated using the standard procedure for SPoly© coatings [10]. The joint members were then adhered together by applying 0.045 g of an equimolar mixture of epoxy, Epon 828®, with methylene dianiline (MDA) curing agent, uniformly on the area of contact and cured at 120°C for 1 hour and at 150°C for 2 hours. The preparation of joints for the torsional tests has been described in detail by McCarvill and Bell [14]. For water immersion tests, the coated substrates were soaked in 60°C water for 3 days before testing. The testing was conducted on an Instron Universal Testing Machine (Model TM-S) at room temperature and humidity (lab conditions). A 4448 N load cell was used to measure the force at break. The crosshead displacement speed was 0.5 cm/min for all the tests. 3. RESULTS AND DISCUSSION

3.1. Effect of HF etching on various alloys The HF treatment was used with five alloys: 2024, 3105, 5000, 6061 and 7075. The compositions of the aluminum alloys are given in Table 1. AA 2024 and 7075 alloys were Al-clad, and that possibly was a reason that 2.5% HF was not an effective pretreatment for these two for the coating process. While the other alloys could be coated after etching in 2.5% HF, AA 2024 and 7075 alloy samples failed to coat after the treatment. Etching with 5% HF solution was found to be effective for all the alloys. Since fluoride ions are well known to induce an increase in the oxide dissolution rate in aluminum [15], a higher concentration HF (also lower pH) aids the process. The alloys, however, exhibited different coating rates; a comparison of the rates is shown in Figure 1. The coating rates presented in Figure 1 are derived from samples which were not wiped after HF etch. As shown, the order of alloys with respect to the coating rate is: 7075, 3105, 5000 > 6061 > 2024. Figure 2 presents the weight loss experienced by the alloys during HF etch. The weight loss for alloys 7075, 5000 and 3105 was higher than for 2024 and 6061. The difference among the etching rates of alloys was likely to have resulted from Table 1. Composition of the Al alloys used for the study Aluminum alloy 2024 3105 5000 6061 7075

Alloy elements, % (aluminum and impurities constitute remainder) Si

0.6

Cu

Mn

Mg

Cr

Zn

4.4

0.6 0.6 0–0.7

1.5 0.5 2.5–4.5 1.0 2.5

0–0.25 0.20 0.25

5.6

0.25 1.6

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Figure 1. Coating rates for various Al alloys (non-wiped) after 1 min etch in 5% HF solution.

Figure 2. Weight loss of the alloys during 1 min etch in 5% HF solution.

the differences in the oxides present on various alloys. Different oxides would be expected to dissolve at different rates in the HF bath. In general, it appeared that a higher weight loss during etching (more etching) led to a higher coating rate. However, the variation in the coating rates could not be explained completely by the weight loss during etching. An interesting observation was that though both


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Figure 3. ESEM pictures of alloys surfaces after 2 min of etch in 5% HF.

AA 7075 and 2024 were Al-clad, the two still exhibited different etching and coating rates. There was no apparent correlation between the bulk composition and the coating rate on the alloys. It is likely that other factors such as surface composition and morphology resulting from the HF etch also affected the rate of coating. Since a detailed surface analysis was beyond the scope of the project, the work to relate the surface composition of the alloys to their coating rate was not pursued. As mentioned above, the alloy samples used in this comparison were not wiped after etch. Some of the alloys, 6061 and 3105 in particular, had a layer of black

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Figure 4. ESEM pictures of alloys surfaces after 2 min etch in 5% HF.

smut formed on their surfaces. It could be largely removed (visually) by wiping the surface with a wet tissue. The surfaces of alloys were observed under the Environmental Scanning Electron Microscope (ESEM) before and after etch to study the effects of the HF etch process. The effect of wiping the surface after etch was also studied. The smut was observed to be a collection of crystalline solid particles dispersed randomly on the surface. The ESEM pictures of surfaces of various alloys (non-wiped) after 2 min etch in 5% HF, are presented in Figures 3 and 4. As the pictures show, there were significant quantities of smut particles present on 6061 and 3105 alloys. Alloys 7075 and 5000 had a moderate number of these particles, while 2024 showed the smallest buildup. The surface of AA 5000 alloy was affected most by the HF etch; it showed a severely etched surface. The wiped surface of AA 7075 alloy, at a higher magnification, is shown in Figure 3. It had a granular appearance, and thus would be expected to have a high surface area. Observations of such granular appearances with F incorporation on the surface and their effect on practical adhesion have previously been reported [16-19]. 3.2. Effect of the smut The coating rates on wiped samples were compared with the rates on non-wiped samples (wiping usually removed the smut visually, but complete removal cannot be assumed). Figure 5 shows the coating rates on AA 2024. A summary of the coating rates on wiped and non-wiped samples is presented in Table 2. The results comparing the weight loss on wiped and non-wiped samples during their pretreatment are shown in Figures 6 and 7.


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Figure 5. Coating rates on AA 2024 alloy for different etch times in 5% HF. Table 2. Coating rates (10-6 g/cm2/s) on wiped (W) and non-wiped (NW) samples Etch

2024 a

time

NW

W

NW

W

NW

W

NW

W

NW

W

20 sec 2 min 5 min

2.3 3.6 4

1.8 2.2 1.3

3.6 4.6 5.1

2.8 0 1.9

3.6 4.9 5.0

2.8 3.1 1.2

3.5 4.3 4.7

2.1 0 0

3.4 4.1 5.9

2.5 2.1 2.1

a b

6061 a

7075 a

3105 b

5000 b

- 5% HF etch - 2.5% HF etch

It was observed that the samples which had smut (non-wiped) exhibited higher coating rates than those without the smut (wiped). The data in Table 2 indicate that this difference was observed with all the alloys. The higher coating rates on the non-wiped samples suggest that presence of smut somehow helps the polymerization on the metal surface. The rate of coating on the non-wiped samples increased with etching time, likely because of an increasing amount of smut. The weight loss results on AA 7075 and 6061 alloys are presented in Figures 6 and 7, respectively. The weight loss for the samples with smut was lower than for samples without it. The difference in the weight losses is a measure of the amount of smut which was wiped away. The smut measurably contributed to the weight of the non-wiped samples. Clearly much more smut was present on AA 6061 than on AA 7075 alloy. More smut was formed at longer etch times and this led to higher coating rates. Although alloys 7075, 2024 and 5000 did not form as much smut as

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Figure 6. Weight loss for wiped and non-wiped AA 7075 alloy samples for different etch times in 5% HF.

Figure 7. Weight loss for wiped and non-wiped AA 6061 alloy samples for different etch times in 5% HF.


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Figure 8. ESEM picture (x2500) of AA 6061 alloy, wiped after 5 min of etch in 5% HF.

6061 and 3105, the difference between coating rates of the wiped and non-wiped samples was still significant, suggesting that even a small amount of smut can make a marked difference. No trend was observed among the coating rates on wiped samples. One possible reason could be the inherent inconsistency in the sample wiping technique. Wiping is not likely to remove the smut completely from the surface; the variation in the quantity of smut left on the surface after the wiping can explain the results. Some unexpected results were observed with wiped samples when different etching times were used. For most of the wiped samples, the coating rate after a 5 min etch was less than the rate after a 2 min or a 20 sec etch. This was not expected; a better etch resulting from a longer etch time was expected to lead to a higher coating rate. However, that was not observed. One infers that in the absence of smut, higher etching (more surface area) may not necessarily result in a higher coating rate. The AA 6061 samples wiped after 2 min of etch coated poorly while the samples wiped after 20 sec or 5 min of etch were coated easily. The ESEM studies showed that AA 6061 samples etched for 2 min had little or no smut left on them after wiping. On the other hand, AA 6061 samples wiped after 5 min of etch gave a different result. A layer of crystals, which could not be removed by wiping, was observed embedded into the surface (Figure 8). The presence of smut particles on the surface assisted the polymerization on the metal surface.

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The AA 3105 samples wiped after 2 min and 5 min of etch did not coat. The ESEM studies on 3105 samples showed that wiping after etch effectively removed the smut. Here again, a lack of smut particles seemed to have caused the failure to coat. Interestingly, the critical role of smut in facilitating polymerization was observed more with AAs 3105 and 6061, which generated more smut than the other alloys. Smut, as is discussed below, was found to contain metals such as Fe and Cu which favor the coating (S-Poly©) process [10, 11]. During etching, these metals tend to concentrate in the smut. As a result of a longer etch, the surfaces of the alloys are likely to become depleted of these trace metals as they go into the smut. This is one possible explanation why wiped AA 3105 samples could be coated after a 20 sec etch but could not be coated after a longer etch. Wiped samples of AAs 2024, 7075 and 5000 could be coated, but their coating rate was significantly lower than that of non-wiped samples. The difference in the coating rates on wiped and non-wiped samples underlines the importance of the smut in the polymerization process. As mentioned before, wiping may not remove smut completely. Also the tendency of smut to adhere to the surface becomes an important factor as well. In case of a sticky smut, the samples may have some smut left on them even after wiping and give variable results. Similarly, during rinsing and drying steps of the non-wiped samples, some of the smut may shed off the surface (caused by water and air-flow), affecting the rate of coating. The practical adhesion of coatings obtained on AA 6061 substrates after HF etch was measured. A comparison of practical adhesion of coatings on wiped and non-wiped samples is shown in Figure 9. The coating on a non-wiped etched sample had a lower adhesion strength, likely due to the presence of smut on the surface. Smut particles are loosely adhered to the substrate and their presence in the coating-substrate interphase might have adversely affected the bond strength. The failure appeared to have occurred in the coating-metal interphase. To study the effect of water on the coatings, practical adhesion strengths of dry coatings were compared to those for the samples immersed in 60°C distilled water for 3 days. A comparison of bond strengths of an etched (non-wiped) and alumina blasted sample is presented in Figure 10. The effect of water was more severe on the non-wiped etched samples than that on the blasted sample, most likely due to the presence of smut particles on the metal surface. The presence of smut particles likely enhances diffusion of water within the interphase. The water within the interphase forms a weakened layer, resulting in loss of practical adhesion. The effect of water on coatings with smut could also be observed in the form of some dark spots which formed under the coating. 3.3. Study of the smut character The ESEM studies on AA 6061 and 3105 samples revealed that the smut crystals were hexagonal in shape (Figure 11). The size of the crystals (in their lateral di-


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Figure 9. Shear strength of S-Poly coatings on HF etched AA 6061 alloy joints, with and without smut.

Figure 10. Comparison of the bond strength of coatings on AA 6061 alloys: dry vs. those immersed in water at 60°C for 3 days.


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Figure 11. ESEM picture (x3000) of smut particles on AA 6061 alloy, after 5 min etch in 5% HF.

mension), as observed on etched AA 3105 alloy, increased with etch time. The smut collected from AA 6061 alloy was analyzed by wide angle XRD. The smut was identified as primarily tri-hydrated aluminum fluoride (AlF3 · 3H2O). The etched surface of AA 6061 alloy was analyzed by means of both XPS and AES. The analyses detected trace quantities of Fe and Cu along with high concentration of Al, F and O on the etched (wiped) surface. The XPS and AES results are summarized in Tables 3 and 4, respectively. Table 3. XPS analyses of etched Al 6061 samples Sample (Al 6061) As-received 5 min etch, non-wiped 5 min etch, wiped

% of the main elements detected on the surfaces of the samples Al

O

C

F

5.9 17.5

48.1 16.6

28.5 10.2

55.1

24.8

41.1

21.1

10

Mg

Si

12.9

2.8

Cu

Fe

Zn 0.5

1.0

0.5

0.3


79

Table 4. The elements detected by AES on the HF etched surface of Al 6061 alloy Sample (Al 6061)

Elements detected on the surface

As-received Non-wiped, 5 min etch Wiped, 5 min etch

Al, O, C, Mg, Ca, Cl Al, O, F, C Al, O, C, Mg, F, Cu, Fe, Zn, Ca

Table 5. ICP-AES results on Al 6061 alloy and the smut collected from the alloy Sample

Al 6061 Smut

% individual metal based on total metal content as 100% Al

Cr

Cu

Mg

Fe

Mn

98.2 95.55

0.20

0.17 0.95

0.98 0.6

0.32 0.95

0.12 1.95

The ICP-AES analysis of the smut from AA 6061 alloy also showed primarily Al with traces of Cu, Fe and Mn (Table 5). The concentrations of these trace metals in the smut were 3-5 times higher than those detected in the AA 6061 bulk. 4. CONCLUSIONS

The results of HF etching were observed to vary among aluminum alloys. The differences were in terms of the weight loss, the extent of coverage with the residual product (smut) and the surface morphology obtained after etching. This resulted, in turn, in the differences in coating rates obtained on the alloys. The smut was observed to have a significant effect on polymerization rate on the metal surface. In some alloys the presence of smut was found to be more critical for the coating formation than in others. The presence of the smut affected the adhesion strength of the coating, and resulted in lowering the adhesion strength further in the presence of water. The smut collected from AA 6061 alloy was found to consist of hexagonal crystals of tri-hydrated aluminum fluoride with traces of other metals such as Fe and Cu. Acknowledgements The authors appreciate the advice of Dr. T.R. Hanlon of United Technologies Corporation in this work.

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REFERENCES 1. S. Wernick, R. Pinner and P.G. Sheasby, The Surface Treatment and Finishing of Aluminum and its Alloys; 5th ed., Vols. I & II., ASM International (1987). 2. J.D. Minford, Handbook of Aluminum Bonding Technology and Data, Marcel Dekker, New York (1993). 3. G.W. Critchlow and D.M. Brewis, Int. J. Adhesion Adhesives, 16, 255-275 (1996). 4. A.D. Yfantis and D.K. Yfantis, ATB Metall., 37 (2-4), 25-30 (1997). 5. F.J. Monteiro and M.A. Barbosa, Surface Coatings Technol. 35, 321-331 (1988). 6. J.W. Golby and J.K. Dennis, Surface Technol. 12, 141-155 (1981). 7. N.T. McDevitt, W.L. Baun and J.S. Solomon, WPAFB technical report, AFML-TR-75-122 (1975). 8. J.S. Solomon and W.L. Baun, in: Surface Contamination-: Genesis, Detection, and Control, K.L. Mittal (Ed.), Vol. 2, pp. 609-634, Plenum Press, New York (1979). 9. W.L. Baun, N.T. McDevitt and J.S. Solomon, paper presented at the Symposium on Surface Analysis of ASTM Committee E-2 on Emission Spectroscopy, Cleveland, Ohio (1975). 10. R. Agarwal and J.P. Bell, Polym.Eng. Sci, 38, 299 (1998). 11. X. Zhang and J.P. Bell, J. Appl. Polym. Sci., 66, 1667 (1997). 12. K. Nainani, Thesis, University of Connecticut, Storrs CT (1998). 13. J.P. Bell and C.J. Lin, J. Appl. Polym. Sci. 16, 1721-1733 (1972). 14. W.T. McCarvill and J.P. Bell J. Adhesion, 6, 185 (1974). 15. A.D. Yfantis and D.K. Yfantis, ATB Metall., 37, 25-30 (1997). 16. W.L Baun, WPAFB technical report, AFML Report TR-79-4172 (1979). 17. W.L Baun, WPAFB technical report, AFML Report TR-79-4178 (1979). 18. W.L Baun, WPAFB technical report, AFML Report TR-79-4155 (1980). 19. W.L. Baun and J.S. Solomon, WPAFB technical report, AFML Report TR-79-4165 (1979).


Surface modification of metals by silanes DANQING ZHU and WIM J. VAN OOIJ∗ Department of Materials Science and Engineering, University of Cincinnati, Cincinnati, OH 45221-0012

Abstract—Various types of silane films on metals were evaluated in a series of ASTM tests for the purpose of corrosion protection. The silane films deposited on AA 2024-T3 were characterized using FTIR-RA. The results showed that: (1) some bis-type silane films, such as bis-[3triethoxysilylpropyl]tetrasulfide (bis-sulfur silane) provided both excellent bare corrosion protection and paint adhesion to different metals; (2) a novel water-borne silane mixture also offered outstanding corrosion protection and paint adhesion to metals, the performance of which was equal to the bis-type silane films; and (3) the structure of the bis-sulfur silane film was affected by factors such as the type of solvent in the silane solution and the nature of the metal substrate on which the silane film was deposited. An interfacial layer at the silane/metal interface was investigated using an FT-IR technique. This interfacial layer is believed to play a key role in the corrosion inhibitory performance of silane films on metals. Keywords: Silanes; Fourier-Transform Infrared Spectroscopy; AA 2024-T3; silane; coatings; corrosion.

1. INTRODUCTION

Silane treatment as a potential replacement for conventional chromating of metals has been studied extensively in our laboratory for several years [1–4]. Recently, bis-type silanes, such as bis-[triethoxysilyl]ethane (BTSE), bis-[trimethoxysilylpropyl]amine (bis-amino silane) and bis-[3-triethoxysilylpropyl]tetrasulfide (bissulfur silane), have become a new focus of our study due to their even better corrosion protection on a wider range of metals or alloys, as compared with monotype silanes. Bis-type silanes had been used primarily as crosslinkers for silane coupling agents until their anticorrosion behavior on a variety of metals or alloys was uncovered [1–6]. Bis-type silanes are usually applied onto metal surfaces from their water/alcohol solutions by a dipping process. The silane-treated metals are then dried in air. The crosslinking of the silane films on the metals in air is rapid, and can be accelerated at elevated temperatures (~ 100 °C). In a silane/metal system it is ∗

To whom all correspondence should be addressed. Phone: 513-556-3194, Fax: 513-556-3773, E-mail: [email protected]

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D. Zhu and W.J. van Ooij

thought [7] that two types of chemical bonds are formed at the interface: (1) siloxane bonds (SiOSi) in the crosslinked silane film; and (2) metallo-siloxane bonds (MeOSi) at the silane/metal interface. Compared with mono-type silanes, bis-type silanes contain more silanol groups if the hydrolysis is complete. Both crosslinking density of the silane film and the number of metallo-siloxane bonds per unit area at the silane/metal interface could be higher. Consequently, an improvement in water resistance and adhesion of the metal to the silane film would be expected, assuming that other functional groups in bis-type silanes do not have detrimental effects. It was found in most cases [2, 3, 6] that bis-type silanes indeed gave a much improved paint adhesion to metal substrates, and, correspondingly, offered better corrosion protection of metals with and without painting. It should be pointed out that although bis-amino silane performs very well when used under paints, it cannot be used alone for the purpose of bare corrosion protection (i.e., without further painting). In a seawater immersion test (i.e., a 3.5% NaCl neutral aqueous solution), for example, AA 2024-T3 alloy coated with the bis-amino silane film applied from a 5% bis-amino water/methanol solution showed some local film delamination in the early stage of the immersion (within 1 day). A complete delamination occurred within 2 hrs of immersion under the harsher condition by adjusting the pH of the NaCl solution to 2, and severe corrosion in the delaminated area was evident soon after. This indicates that the hydrophilic secondary amino groups (-NH-) in the bisamino silane have a detrimental effect on the bare corrosion protection of metals. Subramanian and Van Ooij [2] observed a similar phenomenon with γ-aminopropyltriethoxysilane (γ-APS). They reported that amino groups in γ-APS preferred to bond to metal surfaces, which made the interface between γ-APS silane and metals more hydrophilic, leading to the delamination of the silane film from the substrates more easily by the attack of water. Such delamination, however, has never been observed for the other bis-type silanes. Besides the above bis-type silanes, a novel water-borne silane, a mixture of bisamino and vinyltriacetoxysilane (VTAS), has been developed recently in our laboratory. The major advantages of this silane mixture are listed as follows. (1) The hydrolysis of the mixture is instantaneous. The hydrolysis of VTAS is expressed as: H2C=CH-Si(Ac)3 + 3H2O ⇔ H2C=CH-Si(OH)3 +3HAc

(1)

The acetic acid (HAc) produced from reaction (1) makes the pH of the VTAS aqueous solution very low, e.g., pH 2-3 at a concentration of 5%. The solution at such low pH is unstable, and becomes hazy rapidly due to condensation of the silanol groups. The basic bis-amino silane is therefore mixed with VTAS in certain ratios, functioning as a pH adjuster. (2) The silane solution is almost alcohol-free, i.e., only de-ionized water (DI) water is used as the solvent. Only a small amount of methanol is formed from the rapid hydrolysis of the methoxy ester groups in the bis-amino silane.

Surface modification of metals by silanes

83

(3) The shelf lives of these aqueous solutions at certain mixing ratios are fairly long, more than 4 months. (4) The corrosion performance of this mixed silane film on several alloys equaled that of other bis-type silane films in a series of ASTM industrial tests. In order to obtain a better understanding of the anticorrosion behavior of these silane films, characterization studies were carried out. Fourier-Transform infrared spectroscopy in a reflection absorption mode (FTIR-RA) was employed extensively in this work. The effects of solvent type and metal substrates on the resulting silane film structures were investigated, as well as the interfacial layer suggested in silane/Al system. 2. EXPERIMENTAL

2.1. Materials a) Silanes. Silanes and their structures used in this work are listed in Table 1. Table 1. Silanes and their structures Silane

Structure

Bis-[triethoxysilyl]ethane (BTSE) Bis[3-(triethoxysilyl)propyl]tetrasulfide (Bissulfur silane)

(H5C2O)3Si(CH2)2Si(OC2H5)3 (H5C2O)3Si(CH2)3S4(CH2)3Si(OC2H5)3 (subscript 4 = the average number of sulfur atoms per molecule) (H3CO)3Si(CH2)3NH(CH2)3Si(OCH3)3

Bis(trimethoxysilylpropyl)amine (Bis-amino silane) Vinyltriacetoxysilane (VTAS)

(CH3COO)3SiCH=CH2

The bis-amino and bis-sulfur silanes (Silquest® A1170 and Silquest® A1289, respectively), were provided by Witco Co. (Tarrytown, NY), and VTAS was purchased from Gelest Co. (Tullytown, PA). They were used without further purification. To hydrolyze the bis-sulfur and bis-amino silanes, DI water and/or alcohol were added to the unhydrolyzed silanes. In a 5% bis-sulfur water/ethanol silane solution, for example, the volume ratio between the bis-sulfur silane, DI water and ethanol was 5:5:90. The natural pH of a 5% water/methanol bis-amino silane solution is around 10.8. The solution at such high pH would gel within a few minutes. Acetic acid, therefore, was added to lower the pH of the bis-amino silane solution to 7–8. Preparation of water-soluble silane mixture of bis-amino and VTAS was much simpler than the above silanes. For a 2% solution, 2 parts of the silane mixture of bis-amino and VTAS were simply added into 98 parts of DI water, followed by stirring for a certain period until the solution became clear. Ta-


84

ble 2 shows the solvents commonly used and the pH values of the silane solutions. Table 2. Silanes, solvents and the pHs of the silane solutions Silane

Bis-polysulfur silane Bis-amino silane Mix (Bis-amino silane+VTAS)

Ratio of Solvents

pH

(v/v)

2%

5%

Ethanol/DI water: ~ 90/5 Methanol/DI water: ~ 90/5 DI water only

Natural pH5-6

Natural pH=5-6

pH=7-8 (adjusted by acetic acid) Natural pH=4-5 (at the mixing volume ratio of bis-amino to VTAS of 3:2)

pH=7-8 (adjusted by acetic acid) Natural pH=4 (at the mixing volume ratio of bis-amino to VTAS of 3:2)

Two solution concentrations were used in this study. 2% silane solution was used mainly for improving the paint adhesion to metal substrates and 5% for bare corrosion protection of metal substrates without further painting. 5% solution was also employed in the IR characterization work. b) Metals/alloys. Metals and alloys (15 cm ×10 cm × 0.07 cm) used in this work were as follows: ● 2024 Al alloy in T3 condition (AA 2024-T3, purchased from ACT Laboratories, Hillsdale, MI) ● Cold-Rolled Steel (CRS, purchased from ACT Laboratories, Hillsdale, MI) ● Hot-Dip Galvanized steel (HDG, provided by Bethlehem Steel Co., Bethlehem, PA); the Zn coating thickness was 20 µm. ● Galvalume ® (GVL provided by Bethlehem Steel Co., Bethlehem, PA); the 55% Al-Zn coating thickness was 7 µm. ● Mg alloy (AZ 91B, Norsk Hydro Research Centre, Porsgrunn, Norway) c) Paints. Polyester powder paint from Ferro Corp. and polyurethane powder paint from O'Brien Corp. were used in this study. The paint thicknesses were between 30 to 45 µm and were cured for approximately 15 min at 200 °C. d) Controls. Untreated and chromated panels were used as the controls in ASTM tests. The untreated panels were cleaned with an AC 1055® alkaline cleaner only; the chromated panels were treated with a chromate product with a trade name CHEM-COTE 3509® (Brent International, Lake Bluff, IL).


85

e) Silane treatment of metals. The metals/alloys to be treated were cleaned with a diluted AC1055® (Brent International, Lake Bluff, IL) alkaline cleaner at 65– 75 °C for 3–5 min, followed by rinsing with tap water and drying with compressed air. The cleaned panels were dipped into the silane solutions for 30 s, and then blow-dried with air at room temperature. For the purpose of bare corrosion protection (without further painting), the silane-treated panels were cured at 100 °C for 10 min to obtain an extensively cross-linked structure in the silane films; otherwise the silane-treated panels were not heat-treated. 2.2. Test procedures a) ASTM tests. Three types of ASTM tests [8] were used to evaluate the corrosion performance of silane-treated metals and alloys in this work. ● ASTM B117 (Salt Spray Test): to evaluate the bare corrosion performance of silane-treated metals and alloys. ● ASTM E643-84 (Ball Punch Deformation Test): to evaluate the formability of silane films on metal and alloy sheets. ● ASTM D 1654-92 (Corrosion Test for Painted or Coated Metals and Alloys): to evaluate the corrosion performance of painted alloys with silane pretreatments. b) Characterization of silane films. The structures of silane/metal systems were investigated using FTIR-RA techniques. Experiments were carried out to study: (1) the effect of solvent type on the structure of the bis-sulfur silane film on AA 2024-T3; (2) the effect of metal substrates on the structure of the bis-sulfur silane film; and (3) observation of the interfacial structure in the bis-sulfur silane coated AA 2024-T3 system. FTIR-RA measurements were mainly conducted in the midIR range (e.g., 4000–400 cm–1) on a BIO-RAD infrared spectrometer. The resolution was 4 cm–1, and the number of scans was 256. An FTIR study was also done on a Nicolet® IR spectrometer in the far-IR range from 1800 to 150 cm–1, in order to investigate the interfacial layer in silane-treated metal systems. 3. RESULTS AND DISCUSSION

3.1. ASTM test results 3.1.1. Bare corrosion protection of silane-treated alloys Some bis-type silane films protect a variety of metal substrates against corrosion without topcoats. Empirically, it was found that a thickness of silane films above 200 nm was required for bare corrosion protection. The thickness values of silane films applied onto metals can be controlled by the silane solution concentration. This is because a linear relationship was consistently observed between the film thickness as measured by ellipsometry and the corresponding silane solution con-

86


centrations [1, 6]. The optimum concentration for bare corrosion protection is about 5% for all silanes reported [1]. Figure 1 shows AA 2024-T3 panels after 7 days (168 hrs) of salt spray testing. It can be seen that the silane-treated AA 2024-T3 panels retained their original shiny surface with only slight edge corrosion after the test, while the untreated AA 2024-T3 corroded severely and the chromated alloy surface also showed slightly more corrosion than the silane-treated one. Figures 2 and 3, respectively, present bis-sulfur silane-treated GVL and HDG panels after the deformation test followed by the salt spray test according to ASTM E643-84 and ASTM B117. The aim of this work was to test the bare corrosion behavior of the deformed bis-sulfur silane films on alloys. Apparently, the deformed areas on both GVL® and HDG panels (i.e., the domes in the center of the panels) were protected by the silane films in comparison to the untreated panels. The silane-treated deformed area on the GVL® shows no corrosion; the silane-treated deformed dome on HDG has a few small pits. 3.1.2. Corrosion protection of silane-treated alloys under paints When silanes are used in the pretreatment before painting, the optimum concentration of silane solutions in this application normally is 2% [1]. Figures 4 and 5, respectively, present polyurethane (PU) and polyester (PET) powder-painted AA 2024-T3 panels after 1008 hours of salt spray test (SST). The panels were subjected to different pretreatments (i.e., silane treatment and chromating) before painting and corrosion testing. The silane used here was a water-soluble mixture of bis-amino silane and VTAS. It is evident in Figures 4 and 5 that both silane treatment and chromating provide equally good paint adhesion and corrosion protection for the PET and PU painted AA 2024-T3 panels. The untreated AA 2024-

Figure 1. AA 2024-T3 panels after 7 days (168 hrs) of Salt Spray Test (SST); (a) untreated, alkaline cleaned only; (b) chromated (CHEM-COTE 3509); and (c) mix (bis-amino silane + VTAS) pretreated, (Bis-amino/VTAS=1.5/1(v/v), 5%, natural pH 3.8).


87

T3, on the contrary, shows a large degree of delamination with the PU paint and a complete delamination from the PET paint. PET powder-painted GVL panels were also subjected to salt spray test (SST), as shown in Figure 6. The silane treated GVL panel presents a slight local paint delamination at one end of the scribe line on the surface after 1008 hours of SST exposure, whereas the blank panel (i.e., untreated) totally failed in the test, showing a complete paint delamination and severe corrosion of the substrate nearby the scribe line.

Figure 2. GVL panels after deformation test followed by SST for 48 hrs; (a) untreated, alkaline cleaned only; (b) bis-sulfur silane treated, (5%, natural pH 5-6).

Figure 3. HDG panels after deformation test followed by SST for 48 hrs; (a) untreated, alkaline cleaned only; (b) bis-sulfur silane treated, (5%, natural pH 5-6).

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Figure 4. Polyurethane powder painted AA 2024-T3 panels after 1008 hrs of SST; (a) untreated, alkaline cleaned only; (b) chromated (CHEM-COTE 3509); and (c) mix (bis-amino silane + VTAS) pretreated, (bis-amino silane/VTAS=1.5/1(v/v), 2%, natural pH 4-5).

Figure 5. Polyester powder painted AA 2024-T3 panels after 1008 hrs of SST; (a) untreated, alkaline cleaned only; (b) chromated (CHEM-COTE 3509); and (c) mix (bis-amino silane + VTAS) pretreated, (bis-amino silane/VTAS=1.5/1(v/v), 2%, natural pH 4-5).

3.2. Characterization of silane films on metals 3.2.1. The effect of organic solvents on the film structure of bis-sulfur silane deposited on AA 2024-T3 It has been reported [9] that the hydrolysis of silanes in protic solvents, such as water and ethanol, approaches equilibrium rather than completion, due to the competition between silane hydrolysis and alcoholysis. The hydrolysis process in non-protic solvents such as dioxane, on the contrary, tends to completion. In the case of the bis-sulfur silane, we also observed a similar phenomenon [5, 10]. In


89

Figure 6. Polyester powder painted GVL panels after 1008 hrs of SST; (a) untreated, alkaline cleaned only; and (b) bis-sulfur silane pretreated, (2%, pH 5-6).

this work, one critical issue that needed to be clarified is whether or not the solvent type would further affect the corrosion performance of the resulting bissulfur silane films through influencing the hydrolysis process in the silane solution. To clarify this, FTIR-RA and salt immersion test were employed here. Two types of solvents studied here were ethanol and dioxane. The former is a commonly-used solvent for the bis-sulfur silane solution. Two kinds of bis-sulfur silane films were deposited on AA 2024-T3 from 5% ethanol-based and dioxanebased bis-sulfur silane solutions, followed by curing at 100 °C for 10 min, to make the films sufficiently crosslinked. The FTIR-RA measurements were then conducted on these two silane-coated AA 2024-T3 specimens, and the results are shown in Figure 7. The IR peak assignments are given in Table 3. In Figure 7, a broad band at around 3300 cm–1 appears for both silane films, indicating that there still exists a considerable amount of hydrogen-bonded SiOH groups in the films after curing at 100 °C for 10 min [10]. Two peaks, 2974 cm–1 and 953 cm–1, correspond to the undydrolyzed ester groups (-SiOC2H5). The one at 2974 cm–1 is due to the CH stretching vibration in the ester groups, and the other at 953 cm–1 is attributed to the SiO deformation bending mode in the ester groups. It is noted that for the ethanol-based bis-sulfur silane film these two peaks are pronounced. In the case of the dioxane-based bis-sulfur silane film, however, the peak at 2974 cm–1 has totally disappeared and that at 953 cm–1 has diminished. This difference indicates that the ester groups have been converted completely to silanols in the dioxane-based solution but not in the ethanol-based solution. In other words, a complete hydrolysis of the bis-sulfur silane can be obtained using dioxane as a solvent instead of ethanol. As also shown in Figure 7, only one broad and strong band is seen for the ethanol-based silane film at around 1130 cm–1, which corre-

90


Figure 7. FTIR-RA spectra of ethanol- and dioxane-based bis-sulfur silane films on AA 2024-T3. The silane-treated specimens were cured at 100 °C for 10 min. before IR measurements; the concentration of bis-sulfur silane solutions was 5% with a natural pH 5-6; IR measurements were conducted on a BIO-RAD Win-IR spectrometer. Incident angle: 75°, resolution: 4 cm-1, and number of scans: 100. Table 3. Assignments of absorption bands for bis-sulfur silane films on AA 2024-T3 [11, 12] Absorption band (cm-1)

Assignment

3336, 3367 2974, 2924, 2886 1130, 1030 953 886

Si-OH (H-bonded) Si-O-C2H5 (C-H str.) Si-O-Si (Si-O asym.) Si-O-C (Si-O asym.) Si-OH (H-bonded)

sponds to SiOSi bonds. In the case the dioxane-based film, however, two intense peaks are observed at 1130 cm–1 and 1030 cm–1 for SiOSi bonds. The appearance of these double peaks for SiOSi bonds has been reported to be associated with long SiOSi chains in silane films [11, 12], which apparently can be formed in completely hydrolyzed films and not in partially hydrolyzed films. These results allow a conclusion that the type of organic solvent does affect the structures of the resulting bis-sulfur silane films. Figures 8a and b illustrate schematically the molecular structures of these two types of bis-sulfur silane films on the Al alloy, on the basis of the structural information extracted from the IR spectra discussed above. In the ethanol-based bis-


91

sulfur silane film (Figure 8a), there still exists a large number of unhydrolyzed ester groups (OR) attached to Si and hydrogen-bonded silanols groups (SiOH). These unhydrolyzed ester groups would further hydrolyze to silanol groups in the existing film when in contact with water or moisture [10]. The silanol groups generated in the interfacial region might react with the oxide layer on the Al alloy surface to form an interfacial layer enriched with AlOSi bonds and those in the outer film might condense with themselves to form an extensively crosslinked three-dimensional siloxane network [10]. The structure of the dioxane-based bissulfur silane film on the Al alloy is given in Figure 8b. No unhydrolyzed ester groups are shown in the film, as the hydrolysis in the dioxane-based silane solution tends to completion. The silanol groups in the film can condense readily and completely provided that the curing time is long enough. It seems that the SiOH groups condense in the film to form long siloxane chains, as was supported by the appearance of the double peaks at 1130 cm-1 and at 1030 cm–1. The corrosion performance of these types of bis-sulfur silane films on AA 2024-T3 was evaluated by a salt immersion test (i.e., 3.5%NaCl aqueous solution, pH 7). After 7 days of immersion, no significant difference was seen between these two types of silane films (Figure 9). This indicates that different organic solvents do not have a noticeable effect on the corrosion performance, although they do affect the film structures by influencing the hydrolysis of the bis-sulfur silane in different ways. Based on this result, another interesting conclusion can be drawn here that a completely hydrolyzed silane solution is not necessary for optimum corrosion protection of AA 2024-T3. 3.2.2. The effect of metal substrates on the film structure of bis-sulfur silane In practice, it is often observed that the same silane film performs differently when deposited on different metal substrates. So in order to understand this phenomenon we characterized the bis-sulfur silane films deposited on four different metal substrates. The following common metals/alloys were selected here: AZ 91B (Mg alloy), CRS (Fe alloy), AA 2024-T3 (Al alloy) and HDG (Zn-coated steel). All the metals were treated with the same 5% bis-sulfur silane water/ethanol solution, followed by curing at 100 °C for 10 min before FTIR-RA measurements. The results are shown in Figure 10, and the corresponding IR peak assignments are similar to those given in Table 3. In the SiOH region (3500–3000 cm–1 in Figure 10), a broad band is seen around 3300 cm–1, indicating that H-bonded SiOH groups still exist in the bis-sulfur silane films formed on all the alloys. The intensity of the band, however, depends strongly on the metal substrate. The intensity in the spectrum of the film formed on AZ 91B (Mg alloy) is the smallest; while the intensity in the spectrum obtained from HDG (Zn-coated steel) is the greatest. Obviously, the film on the AZ 91B has been almost fully crosslinked after cured at 100 °C for 10 min, as the very broad band at 3392 cm–1 has diminished compared with the others. The film on HDG, on the contrary, is the least crosslinked because a considerably intense broad band is seen at 3350 cm–1. Further, in the siloxane range (1200–1000 cm–1

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Figure 8. Schematic structure of the two types of bis-sulfur silane films on Al alloy; (a) ethanolbased bis-sulfur silane film; and (b) dioxane-based silane film.

in Figure 10), only one broad and intense band around 1130 cm–1 due to SiOSi bonds is observed for the silane films on CRS, AA 2024-T3 and HDG; whereas double peaks at 1030 cm–1 and at 1130 cm–1 are seen for AZ 91B. As discussed earlier, the appearance of these double peaks for SiOSi is associated with long chains of SiOSi bonds [11, 12]. Apparently, the structure of the bis-sulfur silane film deposited on the Mg alloy (AZ 91B) is different from the others. The one on the Mg alloy is characteristic of long chains of SiOSi bonds. This phenomenon may be explained by taking into account the highly basic nature of magnesium oxide (MgO).


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Figure 9. 7-day seawater immersion test on different types of bis-sulfur silane films on AA 2024T3; (a) Ethanol-based bis-sulfur silane film on AA 2024-T3; (b) Dioxane-based bis-sulfur silane film on AA 2024-T3.

Figure 10. FTIR-RA spectra of ethanol-based bis-sulfur silane films on different metals; (a) HDG; (b) AA 2024-T3; (c) CRS; and (d) AZ 91B. The silane-treated specimens were cured at 100 °C for 10 min. before IR measurements; the concentration of bis-sulfur silane solutions was 5% with a natural pH 5-6.

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It is known that the equilibrium pH of MgO is ~ 11, i.e., very basic in nature [13]. At pH < 11, MgO undergoes acidic dissolution according to the reaction, MgO + H2O ⇔ Mg2+ +2OH–

(2) 2+

It is noted that dissolution of MgO not only generates magnesium ions (Mg ), but also produces hydroxide ions (OH–). In this work, the pH of the 5% water/ethanol bis-sulfur silane solution was 5–6, thus some of the MgO layer was very likely to dissolve during the process of silane treatment, resulting in a high pH region close to the metal surface, as shown schematically in Figure 11a. Both hydrolysis and condensation would be accelerated by a base-catalyzed mechanism in this high pH region. As a result, long SiOSi chains were formed in the resulting silane films, as evidenced by the appearance of the double SiOSi peaks for AZ 91B in Figure 10. Grubb observed a similar effect caused by alkali metal hydroxides [14]. He stated that the condensation of trimethylsilanol in methanol was catalyzed by alkali metal hydroxides. The rate of condensation was dependent on the concentration of hydroxide anion, but not on the alkali metal cation. In contrast with MgO, the other oxides, such as Al2O3, ZnO, and FeO, are much more stable at neutral pH. Therefore, oxide dissolution is not expected on the other alloy surfaces. In the case of AA 2024-T3, for example, a hydroxide layer containing Al hydroxyls (AlOH) is formed (Figure 11b). The as-generated AlOH groups would react with SiOH groups in the silane solution, forming Al-siloxane (AlOSi) later on. Figures 12a and b show schematically the interfacial structure of silane/metal systems, on the basis of the structural information extracted from the FTIR-RA spectra in Figure 10. Figure 12a illustrates that the interfacial structure of the bissulfur sialne treated Mg system is affected by the local alkalization discussed above. The new SiOSi peak at 1033 cm–1 in Figure 10 clearly reflects this pH effect on the structure. As a comparison, Figure 12b depicts the interfacial structure of the silane/Al system. AlOSi covalent bonds form at the silane/Al interface as a result of the condensation between SiOH groups and AlOH groups at the alloy surface. A small amount of H-bonded SiOH and AlOH may also exist at the interface as a consequence of the equilibrium [7]. 3.2.3. Observation of the interfacial layer in silane/metal systems It was found in our recent electrochemical impedance spectroscopy (EIS) work with a non-corrosive electrolyte (0.5M K2SO4) [10] that there existed three time constants in the EIS spectra for a bis-sulfur silane coated AA 2024-T3 system. These three time constants were then assigned to, from outside to inside, the outer silane film, a new interfacial layer between the silane film and the oxide layer, and the inner oxide layer. FTIR-RA measurements also showed some evidence which supported this suggestion [10]. It was further speculated that the interfacial layer was very likely the product of the reaction (1) between AlOH groups on the AA 2024-T3 surface and the SiOH groups in the silane film, and (2) among the SiOH groups in the interfacial region. This layer was, therefore, assumed to be composed of aluminum silicates ([Al2O3•XSiO2]) [10, 15].


95

To further clarify this inorganic interfacial layer, the FTIR-RA spectra of two types of silane films on AA 2024-T3 were obtained in the far-IR range (2000 to 200 cm–1), as shown in Figure 13. The silane films were deposited on AA 2024-

Figure 11. Schematic of the interfaces between metals and bis-sulfur silane water/ethanol solution; (a) MgO/silane-solution; (b) Al2O3/silane-solution.

96


Figure 12. Schematic of the interfacial regions between metals and bis-sulfur siloxane network; (a) Mg /siloxane-network; (b) Al/siloxane-network.

T3 from the 5% silane solutions of bis-amino and bis-sulfur silanes. The silane films were cured at 100 °C for 10 min before the IR measurements. In Figure 13 two broad but weak bands in the region from 1200 cm–1 to 800 cm–1 are observed for the untreated AA 2024-T3, which are associated with the Al oxide layer on the alloy surface. The AA 2024-T3 substrate coated with the bis-amino silane film shows a strong and broad band in the region from 1200 cm–1 to 1000 cm–1, known as the SiOSi region, and a medium one around 500 cm–1. The AA 2024-T3 panel coated with the bis-sulfur silane film exhibits similar results: one intense peak at 1100 cm–1 indicating SiOSi bonds and one medium peak around 500 cm–1. As no peaks were observed for the untreated AA 2024-T3 at 500 cm–1, the peak around 500 cm–1 for both silane films on AA 2024-T3 is thus attributed to the silane films.


97

Figure 13. FTIR-RA spectra of untreated and silane-treated AA 2024-T3; (a) AA 2024-T3 untreated; (b) treated with bis-sulfur silane (5%, natural pH 5-6); and (c) treated with bis-amino silane (5%, pH adjusted to 7 with acetic acid). The silane-treated specimens were cured at 100 °C for 10 min. before IR measurements; IR measurements were conducted on a Nicolet spectrometer.

Since few peaks are normally observed for polymeric materials below 600 cm–1, the peaks at around 500 cm–1 are probably due to an inorganic region formed in the silane film and/or at the silane/metal interface [16]. The authors in Ref. [15] assumed that some bands corresponding to metallosiloxane (MeOSi) bonds also appeared in the region from 1200 cm–1 to 1000 cm–1. It is, however, difficult to extract such information from Figure 13, as the 1200– 1000 cm–1 region is dominated by the bands corresponding to SiOSi bonds. The bis-sulfur silane solution was further diluted from 5% to lower concentrations, in order to deposit thinner films on AA 2024-T3. More specific information on the interfacial region is thus possible to be extracted from the spectra of the thinner films. Figure 14 compares the FTIR-RA spectra of bis-sulfur silane films on AA 2024-T3 alloy applied from the concentrations of 5%, 1%, 0.5%, and 0.1%. Some weak peaks in the case of 5% can be distinguished clearly at low concentrations (1–0.1%) (Figure 14). In the region from 550 cm–1 to 400 cm–1, three medium peaks have become distinct with decreasing concentrations. These peaks may again be associated with the inorganic layer at the silane/metal interface. The peaks could be attributed to the Si-O deformation bending mode for SiOSi in inorganic silica or silicates [16]. We thus tentatively postulate an interfacial structure of the organic/inorganic silicates of the type [Al2O3•XSiO2], in which X is an organic group, such as –(CH2)3, still attached to the Si atom.


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Figure 14. FTIR-RA spectra of ethanol-based bis-sulfur silane films on AA 2024-T3 applied from silane solutions with different concentrations; (a) 0.1%; (b) 0.5%; (c) 1%; and (d) 5%. All specimens were polished and finished with 1 µm alumina paste; the silane-treated specimens were cured at 100 °C for 21 hrs. before the measurements.

4. CONCLUSIONS ●

●

●

Bis-type silanes and a novel waterborne silane mixture, bis-amino and VTAS, provide excellent bare corrosion protection of and paint adhesion to a variety of alloys, as shown in a series of industrial ASTM tests. The type of organic solvent in the bis-sulfur silane solution has a remarkable effect on the hydrolysis of the silane solution. The FTIR-RA results indicated that the hydrolysis of bis-sulfur silane can be accelerated using dioxane as a solvent instead of ethanol. The structure of the bis-sulfur silane films and their corresponding corrosion performance are also influenced by the solvent type in the bis-sulfur silane solution. The ethanol-based bis-sulfur silane film shows only a partly crosslinked structure due to the incomplete hydrolysis in the ethanol-based solution, while that of dioxane-based film has a fully crosslinked structure as a result of the complete hydrolysis in the dioxane-base solution. The difference in the corrosion performance of these two types bis-sulfur silane films on AA 2024-T3, however, was insignificant. This indicates that a complete silane hydrolysis in the silane solution is not necessary for bare corrosion protection of metals.


●

●

99

Metal substrates, on which the bis-sulfur silane film is deposited, have a significant effect on the film structure. In the case of the Mg alloy (AZ 91B), the basic MgO dissolves partly in the bis-sulfur silane water/alcohol solution. The as-generated OH– ions increase the local pH close to the alloy surface; as a result, both hydrolysis and condensation of the bis-sulfur silane are accelerated by a base-catalyzed mechanism. The less basic oxide layers on the other alloys, i.e., AA 2024-T3, HDG, and CRS, only form hydroxide layers in the presence of water or moisture, which do not vary the pH nearby the alloy surfaces. Therefore, the bis-sulfur silane films formed on these alloys are similar in structure. The existence of an inorganic interfacial layer at the silane/metal interface was concluded from FTIR studies in both mid-and far-IR ranges.

Acknowledgment The authors are grateful for the financial support provided by Brent International PLC in Bletchley, UK, the use of salt spray facilities at Brent International in Lake Bluff, IL, and the assistance of Thomas Rose for the far-IR measurements at Chemetall/GmbH in Frankfurt, Germany. REFERENCES 1. W.J. van Ooij and T.F. Child, CHEMTECH, 28, 26 (February 1998). 2. V. Subramanian and W.J. van Ooij, Corrosion 54, 204 (1998). 3. V. Subramanian, Ph.D. Thesis, University of Cincinnati, Department of Materials Science and Engineering (1999). 4. W.J. van Ooij, D. Zhu, Guru Prasad, Senthil K. Jayaseelan, Yuan Fu and Niranjan Teredesai, Surf. Engg. 16, 386 (2000). 5. Th. Van Schaftinghen, M.S. Thesis, Vrije Universiteit Brussel, Department of Materials Science and Electrochemistry, Brussels, Belgium (1999). 6. Guru Prasad Sundararajan, M.S. Thesis, University of Cincinnati, Department of Materials Science and Engineering (2000). 7. E.P. Plueddemann, Silane Coupling Agents, 2nd Edition, Plenum Press, New York (1991). 8. R.F. Allen and N.C. Baldini (Eds), 1998 Annual Book of ASTM Standards, ASTM, West Conshohocken, PA (1998). 9. F.D. Osterholtz and E.R. Pohl, in: Silanes and Other Coupling Agents, K.L. Mittal (Ed.), p. 119, VSP, Utrecht (1992). 10. W.J. van Ooij and D. Zhu, Corrosion, 57, 413 (2001). 11. G. Socrates, Infrared Characteristic Group Frequencies, John Wiley & Sons, New York (1994). 12. M.W. Urban, Vibrational Spectroscopy of Molecules and Macromolecules on Surfaces, John Wiley & Sons, New York (1993). 13. E. Ghali, in: R.W. Revie (Ed.), Uhlig’s Corrosion Handbook, 2nd ed, John Wiley & Sons, New York, p. 793 (2000). 14. W.T. Grubb, J. Am. Chem. Soc. 76, 3408 (1954). 15. A.M. Beccaria and L. Chiaruttini, Corrosion Sci. 41, 885 (1999). 16. Infrared Spectroscopy Committee of the Chicago Society for Paint Technology. Infrared Spectroscopy: Its Use in the Coatings Industry. Federation of Societies for Coatings Technologies, Philadelphia, PA (1969).



Application of X-ray photoelectron spectroscopy in assessing the adsorption of siloxane polymers onto E-glass fibers LEANNE G. BRITCHER, DAVID KEHOE and JANIS G. MATISONS∗ Polymer Science Group, Ian Wark Research Institute Special Research Centre for Particle and Material Interfaces, University of South Australia, Mawson Lakes 5095, South Australia, Australia

Abstract—This paper focuses on the use of X-ray Photoelectron Spectroscopy (XPS) in establishing the adsorption of siloxanes and silanes onto E-glass fibers. When a siloxane or silane is adsorbed onto E-glass fibers, the O(1s) and Si(2p) binding energies in the XPS spectra do not shift or broaden significantly to allow identification of the silane/siloxane coating. We have found that the difference between the binding energies of the O(1s) and Si(2p) photoelectron lines, however, can be used to characterize the surfaces of siloxane coated E-glass fibers, as either resembling a siloxane (good coating) or the E-glass fiber (poor coating). The O(1s)-Si(2p) binding energy difference was also used to assess the adsorption of various functional siloxanes to E-glass fibers as a stringent consecutive washing procedure was applied to the surfaces using a variety of selected solvents. Siloxane coupling agent analogs were also prepared and they appeared to interact with the E-glass fiber surfaces in the same way as the analogous alkylsilane coupling agents. DRIFTS and SEM studies have also been used to support the XPS data. Keywords: Siloxane polymers; XPS; surface modification; E-glass fiber; siloxane adsorption.

1. INTRODUCTION

The XPS is a surface sensitive technique, which is capable of providing chemical information to a depth of only a few nanometers, and has been applied to a wide range of modified surfaces [1]. Although the XPS allows quantification of elements at the surface, the interpretation of such data in terms of the ‘degree of coverage’ or coating ‘thickness’ is the subject of much discussion in the literature [2-6]. There have been a number of studies, which have applied XPS analysis to modified E-glass fibers. In most cases, a semi-quantitative approach has been taken. Fagerholm et al. [7] have used XPS (together with streaming zeta potential and dynamic contact angles) to analyze the aqueous application of a surfactant, ∗

To whom all correspondence should be addressed. Phone: 61 8 8302 3162, Fax: 61 8 8302 3755, E-mail: [email protected]

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L.G. Britcher et al.

alkylphenylpoly(oxyethylene)alcohol onto E-glass fibers. The oxygen elemental concentration, almost 50% in the clean fiber, was reduced to 21-24% by application of 0.2-2% solution of the surfactant. If the entire glass surface was being progressively covered by the surfactant, then oxygen percentage should decrease correspondingly. Fagerholm et al. [7] suggested that preferential adsorption onto some of the glass active sites occurred, followed by progressive multilayer adsorption onto the surface patches. Both streaming zeta potential and dynamic contact angle data support this hypothesis. Wang et al. [8] used XPS and Time-of-Flight Secondary Ion Mass Spectrometry (TOF-SIMS) to examine the interaction of hydrolyzed γ-aminopropyltriethoxysilane onto E-glass surfaces. In particular, they were interested in the effects of varying the pH and of treating the finished fibers with hot water. They determined that a pH of 9 resulted in maximum silane adsorption, and as the temperature of the water extraction increased, more adsorbed silane was removed from the surface. Interestingly, the observed aluminium concentration remained the same, as the amount of adsorbed silane increased, suggesting that either aluminium from the glass must have migrated into the silane layer, or the aluminium sites on glass were not covered. The former proposal was supported by the water extraction results, where the elemental concentration of Al was greater after extraction (~ 14%) than that for untreated glass (~ 8%). The Ca level was negligible, after warm water extraction, yet returned to almost 5% after hot water extraction, suggesting that some of the Ca was incorporated into the layer in a water soluble form, such as Ca(OH)2. Wesson et al. [9] used XPS and Electrolytic Thermodesorption Analysis of Water, Inverse Gas Chromatography & Programmed Thermal Desorption to examine the acid-base characteristics of silane treated E-glass fiber surfaces. Aqueous solutions of γ-aminopropyltriethoxysilane (APTES), γ-chloropropyltrimethoxysilane (CPTMS) and methyl-trimethoxysilane (MTMS) were applied to E-glass fibers, and angle resolved XPS was used to determine the orientation of the adsorbed silanes. For APTES, a small percentage of N in the form of NH3+, was found but was observed only at the largest penetration depth (take-off angle 90°). This indicates that only a small percentage of the amino groups are oriented towards the surface, and do in fact react with it. Similarly, the C elemental concentration decreased with increasing penetration depth, indicating that the propyl groups were also oriented towards the surface. The results for CPTMS were surprising, as both the Cl and C elemental concentrations increased with penetration depth. A significant amount of the silane must be oriented with the chlorine and propyl groups towards the glass surface. The aluminium elemental concentration suggested that the silanes did deposit preferentially on aluminols. Pantano and Wittberg similarly used XPS to examine the aqueous application of a number of silane coupling agents onto E-glass fibers [10] (upon silane treatment Ca and Al surface elemental concentrations were reduced). Pantano and Wittberg [10] suggested that the silanes uniformly covered the glass fiber surface, so the thickness, t, of the adsorbed silane film could be determined by:

Application of X-ray photoelectron spectroscopy

t = - λC ln [1 - IC / IC*] . 2/π

103

(1)

where IC is the measured carbon concentration on a particular silane-treated sample, IC* is the carbon concentration that would be measured in an infinitely thick film of that silane and λC is the mean free path of the C 1s photoelectron. Since calcium is present only in the E-glass fiber, the thickness of the overlayer was also determined by: t = - λCa ln [ICa / ICa*] . 2/π

(2)

where ICa is the measured Ca 2p concentration on a particular treated sample, ICa* is the Ca 2p concentration for a clean E-glass surface and λCa is the mean free path of the Ca 2p photoelectron in the silane overlayer. The factor 2/π in these expressions accounts for the curvature of the fiber, which can increase the effective path length of photoelectrons in the surface layer. Pantano and Wittberg [10] then calculated the silane film thickness using equation (2) for each silane, and found that the calculated result corresponded closely to the molecular length of each silane. Such methods make a number of assumptions in order to simplify the calculations. The uniformity of the coating is very much system dependent, and thus may present a real problem in applying these techniques to certain systems. An alternate way of assessing the “degree of coating” takes advantage of the accurately determined binding energies in XPS measurements. Wagner et al. [11] applied XPS to the analysis of 31 silicon-oxygen compounds, and found that the difference between the O(1s) and Si(2p) photoelectron line energies in inorganic silicon compounds ranged between 429.0-429.6 eV. However, for siloxane polymers, the line energy differences between 429.8-430.1 eV, were observed with an accuracy of 0.1 eV reported for the samples run on a non-monochromated source. This difference between the O(1s) and Si(2p) photoelectron line energies may, therefore, be useful to characterise how a silane or siloxane coating covers silicate surfaces. This method has a number of advantages. It is less affected by variations in adventitious carbon levels; it does not rely upon assumptions regarding the uniformity of the coating; and it can be accurately determined from the narrow scan multiplex spectrum of the sample. Previously we have presented the use of Diffuse Reflectance Infrared Fourier Transform Spectroscopy (DRIFTS) to monitor the adsorption of siloxanes onto Eglass fibers [12]. The objectives of this work were to determine if the O(1s)Si(2p) binding energy difference could be used to distinguish siloxanes on E-glass fiber surfaces. This paper discusses the analysis of functionalised siloxane polymers and silanes on E-glass fibers by using the O(1s)-Si(2p) binding energy difference in the XPS spectra.

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2. EXPERIMENTAL SECTION

2.1. Materials Untreated E-glass fibers (water washed only) were obtained from ACI (Australia) (composition: 55% SiO2; 21.5% CaO; 14.5% Al2O3; 6.0% B2O3; 0.6% MgO; 0.9% Na2O and K2O; 0.4% Fe2O3 and 0.6% F2) and used without any further modification. Methylene chloride was dried with anhydrous CaCl2 for 24 hours and then distilled (Ace Chemicals, South Australia). Toluene (Ace Chemicals, South Australia) was distilled over sodium wire with benzophenone until a violet colour developed, indicating the formation of a ketyl radical. Hexane (Shell Australia), was distilled and stored over molecular sieves. Acetone (Chemplex), and AR grade ethanol (Merck) were used as supplied. The following functionalized siloxane polymers were prepared by Duplock et al. [13], from the direct hydrosilation of various poly(dimethylsiloxane-comethylhydrogensiloxane) fluids with the appropriate unsaturated reagents: poly[methyl(3-methacrylatepropyl)siloxane] fluids (degree of polymerization, D.P. = 22 and 33); poly[dimethylsiloxane-co-methyl(ketenimine)siloxane] (D.P. = 20; m = 16, n = 4); α,ω-[(3-amino-n-propyl)terminated-dimethylsilylpoly(dimethylsiloxane)] fluid (D.P. = 19 or 836) and poly[dimethylsiloxane-co-methyl(3aminopropyl)siloxane] (D.P. = 20; m = 16, n = 4). The synthesis of the siloxane couplings agents via hydrosilation of vinyltrimethoxysilane and vinyltris(2methoxyethoxysilane) onto poly(methylhydrogensiloxane) and poly(dimethylsiloxane-co-methylhydrogensiloxane) fluids has been outlined previously [14]. 2.2. Application procedures The application procedure consisted of dissolving the silanes or functionalized siloxane polymers in an appropriate organic solvent (see Tables 1 and 2), and placing this solution (~ 150 cm3) into a 250 cm3 glass vessel together with unsized Eglass fibers (~ 1 g). The glass fibers were left in this solution overnight (16 hours). The fibers were then removed from the solution, and washed with a series of organic solvents (see Table 2) and dried in an oven (3 hours at 100ºC in air). All treated fibers were analyzed by XPS, DRIFTS and SEM. 2.3. Diffuse reflectance infrared Fourier transform (DRIFT) spectroscopy DRIFT spectra were measured at room temperature with a single beam BIO-RAD Model FTS 65 spectrometer in the wavenumber region from 4000 to 400 cm–1, at a resolution of 4 cm–1, using an MCTA detector. The interferogram was apodized using the boxcar method installed in the FTIR software (BIO-RAD, 3200 Data Stations Users Manual, 0910260E, 1989). Signal-to-noise ratio was generally better than 100 to 1, using 256 scans.


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Table 1. Functionalized siloxanes applied to E-glass fibers Functionalized siloxane

Designation CAA1

CAA2

CAA3

CAA4

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Table 1. (Continued) Functionalized siloxane

Designation methacryl 1 (n = 22) methacryl 3 (n = 33)

amino 1 (n = 19) amino 2 (n = 836)

amino 3

imino


107

Table 2. Application procedures for the silanes and functionalized siloxanes Sample name

Siloxane or silane

Conc. (%)

Application procedure

Post-treatment (organic solvent wash)

VTMES1

VTMES

4.98

ethanol

VTMES2

VTMES

4.98

ethanol

VTMS

VTMS

3.40

ethanol

CAA1

CAA1

3.06

toluene

CAA2

CAA2

3.06

toluene

CAA3

CAA3

3.06

toluene

CAA4

CAA4

3.06

toluene

ethanol, methylene chloride (2x75 cm3 of each). methylene chloride, toluene, hexane, acetone, methylene chloride (100 cm3 of each). methylene chloride, toluene, hexane, acetone, methylene chloride (100 cm3 of each). methylene chloride, toluene, hexane, acetone, methylene chloride (100 cm3 of each). methylene chloride, toluene, hexane, acetone, methylene chloride (100 cm3 of each). methylene chloride, toluene, hexane, acetone, methylene chloride (100 cm3 of each). methylene chloride, toluene, hexane, acetone, methylene chloride (100 cm3 of each).

methacryl 1

methacryl 1

3.73

toluene

toluene (2 × 75 cm3)

methacryl 2

methacryl 1

3.73


methacryl 3

methacryl 3

1.41

methacryl 1 + toluene + 0.7% ethylene diamine toluene

methacryl 4

methacryl 3

1.41


amino 1

amino 1

1.70

methacryl 3 + toluene + 0.8% ethylene diamine toluene

amino 2

amino 2

5.20

toluene

toluene (2 × 75 cm3), methylene chloride (3 × 75 cm3)

amino 3

amino 3

4.10

toluene


imino

imino

0.94

toluene




The fibers were mounted parallel to each other in a specially constructed sample holder [12], which was placed in a SpectraTech diffuse reflectance apparatus. Each glass fiber sample was scanned 256 times, with the fibers at an angle of 90°, with respect to the direction of the infrared beam, for maximum signal-to-noise

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Figure 1. DRIFT spectrum of a clean E-glass fiber a) and DRIFT spectra of coated E-glass fiber samples b) VTMES 1 and c) VTMES 2.

ratio. The background spectrum was taken from high purity (IR) grade KBr powder (Merck), placed in a sample cup, and levelled to the top of the cup using a spatula. No pressure was applied to the KBr powder in the cup. DRIFT spectra, such as those in Figure 1, were obtained by subtracting the B-O overtone at 2671 cm–1 of the E-glass fibers from the same overtone on the treated glass fibers, using the BIO-RAD software. 2.4. X-ray photoelectron spectroscopy (XPS) XPS spectra were obtained using a Perkin-Elmer PHI 5600 XPS system with a concentric hemispherical analyzer and Mg Kα non-monchromated X-ray source operating at 300 W. Pass energies of 89 and 18 eV were used for survey and elemental spectra, respectively. Atomic concentrations were obtained using standard sensitivity factors [15]. The angle between the sample surface and the analyzer was fixed at 45°. The pressure during analysis ranged from 10–8 to 10–9 Torr. A five minute Argon ion etch was applied after the initial XPS analysis; the sputter rate was calibrated from the complete removal of a Ta2O5 film on tantalum (5 nm thick). All binding energies (B.E.s) were referenced to the C1s neutral carbon peak at 284.6 eV, to compensate for the effects of surface charging.


109

2.5. Scanning electron microscopy (SEM) A Cambridge stereoscan 100 SEM was employed in this study to routinely examine the modified E-glass fiber surfaces. The treated fibers were coated with a thin (~ 20 nm) evaporated carbon layer to reduce the effects of charging. 3. RESULTS AND DISCUSSION

3.1. Adsorption of siloxane and silane coupling agents When assessing the adsorption of silane and siloxane coupling agents onto Eglass fibers it is difficult to determine the extent of coverage by the binding energies and elemental concentrations in the XPS spectra. Even if there are present other elements besides oxygen, silicon and carbon, it is not possible to determine if the coating is ‘patchy’ or extensively covers the surface. As discussed earlier, Wagner et al. [11] found that the difference between the O(1s) and Si(2p) binding energies could be used to distinguish between silicate and siloxane compounds. This difference maybe useful in characterising the surfaces of silane/siloxane coated E-glass fibers as either resembling that of a siloxane (indicating extensive coverage), or an inorganic silicate (indicating ‘patchy’ coverage). We have been investigating the usefulness of the O(1s)-Si(2p) binding energy difference on a number of different silane/siloxane treated E-glass fibers. When the binding energy difference is ≥429.8 eV, the glass is thought to have mainly a silicone coating; a difference of 429.7 eV is interpreted as being in between a silicone and silicate, while a difference of ≤429.6 eV is interpreted as a ‘patchy’ coating or not detected (based on SEM results). In order to determine the error involved in these measurements we have done triplicate measurements on treated Eglass fibers. Three XPS spectra were obtained on three different days with a fresh sample of the fiber taken for each measurement. The binding energy values did not vary more than 0.2 eV while the O(1s)-Si(2p) binding energy difference error value was 5 µm long) between adjacent fibers on the otherwise generally smooth surfaces of CAA1, CAA2 and CAA3. These SEM micrographs confirm the DRIFTS data suggesting that some siloxane is on the surface, but only rare patches are present, and therefore the O(1s)-Si(2p) binding energy difference shows a predominantly silicate rather than siloxane surface. The SEM micrographs of the E-glass fibers treated with CAA4 are displayed in Figures 6a and 6b. The surface morphology consists of both small (5 µm long) also between adjacent fibers. These micrographs indicate not only that substantial siloxane is on the fiber surfaces (confirming DRIFTS and XPS data), but the coat-

Figure 5. SEM micrographs of coated E-glass fiber samples a) CAA1 b) CAA2 and c) CAA3.

114

Figure 5. (Continued).



115

Figure 6. SEM micrographs of coated E-glass fiber sample CAA4; the top (a) micrograph shows bridges formed between the fibers, while the bottom (b) micrograph shows the siloxane encapsulating the fibers.

116


ing is not uniform (hence the bubbles on the surface). CAA4 not only encapsulates the fiber surfaces, but also can link adjacent fibers through siloxane bridges. XPS, DRIFTS and SEM have shown that the nature of the siloxane backbone plays a significant role in the adsorption of coupling agent analogs onto E-glass fibers. 3.2. Adsorption of methacryl functionalised siloxane polymers Two different methacryl functionalized siloxane polymers were prepared and applied to E-glass fibers from toluene (methacryl 1, D.P. = 22 and methacryl 3, D.P. = 33). Ethylene diamine was also added to toluene solutions of methacryl 1 and methacryl 3, generating two further samples methacryl 2 and methacryl 4 respectively (see also Table 2). Blitz and coworkers [19] investigated the interactions of silanes and silica in the presence of various amines. They concluded that all amines catalyzed the interaction between silanes and the surface hydroxyl groups on silica. The ethylene diamine was added to determine the effect of increasing the pH on the adsorption of the methacryl siloxane polymers. Figure 7 displays the the O(1s)-Si(2p) binding energy differences for methacryl functionalized siloxanes on E-glass fibers. Samples methacryl 1 and methacryl 3 have binding energy differences typical of a siloxane polymer surface. The binding energy differences for methacryl 2 and methacryl 4 indicate that the presence of the ethylene diamine (see Table 2) has not enhanced but rather reduced the ad-

Figure 7. O(1s)-Si(2p) binding energy differences for methacryl functionalized siloxane polymers on E-glass fibers. Where = binding energy (B.E.) difference before a 5 minute Argon etch and = binding energy difference after the samples were etched for 5 minutes by Argon.


117

sorption of the methacryl functionalized siloxanes. DRIFTS and SEM results, presented previously [12], did not indicate clearly whether the amine competed with the methacryl functionalized siloxane for adsorption sites. XPS has now shown that the presence of amine does interfere with adsorption as the O(1s)Si(2p) binding energy differences are lower for samples containing the amine (methacryl 2 and methacryl 4). 3.3. Adsorption of imino and amino functionalized siloxanes We have observed that amino acid functionalized siloxanes groups can adsorb effectively onto E-glass fibers [20] and there has been much discussion in the literature regarding the interaction of such functional groups attached to trialkoxysilanes with glass surfaces [21-22]. DRIFTS and SEM results, presented previously [12], did show adsorption of the siloxane polymers, but they were not able to determine if the siloxane polymers entirely covered the surface or were only present as patches which were observed in the SEM micrographs. Several siloxane polymers bearing amino or imino functional groups were prepared and applied to unsized E-glass fibers, in order to investigate how such functional groups interacted with glass surfaces. An imino functionalized siloxane based on a poly(dimethylsiloxane-co-methylhydrogensiloxane) polymer, with an average of 4 imino groups per molecule was prepared. Both amino 1 and amino 2 samples are aminopropyl-terminated poly(dimethylsiloxane) polymers of different molecular weights (amino 1 D.P. = 19; amino 2 D.P. = 836). Sample amino 3 is a pendant aminopropyl functionalized siloxane, with an average of 4 pendant amino groups per molecule. For the imino functionalized siloxane was prepared by reaction of amino 3 with a ketone. The O(1s)-Si(2p) binding energy differences for E-glass fibers coated with the imino and amino functionalized siloxanes are displayed in Figure 8. All of the amino siloxane samples gave O(1s)-Si(2p) binding energy differences typical for a siloxane surface. For the imino siloxane coated E-glass binding energy difference is 429.7 eV. Clearly the amino groups have adsorbed more effectively onto the glass fiber surface than the imino siloxane and are not easily removed after washing with toluene. 4. CONCLUSIONS

The O(1s)-Si(2p) binding energy difference obtained from the XPS spectra is a useful way to assess the surface coverage of a siloxane on E-glass fibers quickly and effectively. When the XPS data are combined with DRIFTS and SEM results, we have a powerful analytical technique for characterising the adsorption processes of siloxane polymers, bearing diverse functional groups, onto E-glass fibers.

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Figure 8. The O(1s)-Si(2p) binding energy differences for imino and amino functionalized siloxane polymers on E-glass fibers. Where = binding energy (B.E.) difference before a 5 minute Argon etch and = binding energy difference after the samples were etched for 5 minutes by Argon.

Vinylsilanes do absorb onto E-glass fibers from solution but may be effectively removed if the fibers are subsequently washed with a series of appropriate solvents. When the vinylsilanes are attached to a siloxane polymer via hydrosilation, the siloxane backbone determines the effective adsorption onto E-glass fibers. Siloxane polymers bearing amino and methacryl groups remain on the E-glass fiber surface even after washing with toluene, indicating good adsorption onto the Eglass fiber surface. REFERENCES 1. D. Briggs and M.P. Seah (Eds), Practical Surface Analysis by Auger and X-ray Photoelectron Spectroscopy, Wiley, New York (1983). 2. T.L. Barr, Critical Rev. Anal. Chem. 22, 569 (1991). 3. O.A. Baschenko, J. Electron Spect. Rel. Phenom. 57, 297 (1991). 4. L.S. Johansson, Surf. Interf. Anal. 17, 663 (1991). 5. J.E. Fulghum, Surf. Interf. Anal. 20, 161 (1993). 6. K.I. Nishimori and K. Tanaka, J. Vac. Sci. Technol. A8, 3300 (1990). 7. H. Fagerholm, C. Lindsjö, J. Rosenholm and K. Rökman, Colloids Surfaces 69, 79 (1992). 8. D. Wang, F.R. Jones and P. Denison, J. Adhesion Sci. Technol. 6, 79 (1992). 9. S.P. Wesson, J.S. Jen and G.M. Nishioka, J. Adhesion Sci. Technol. 6, 151 (1992). 10. C.G. Pantano and T.N. Wittberg, Surf. Interface Anal. 15, 498 (1990). 11. C.D. Wagner, D.E. Passoja, H.F. Hillery, T.G. Kinisky, H.A. Six, W.T. Jansen and J.A. Taylor, J. Vac. Sci. Technol., 21, 933 (1982).


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12. L.G. Britcher, D.C. Kehoe and J.G. Matisons, Silanes and Other Coupling Agents, Vol 2, K.L. Mittal (Ed.), pp. 99-114, VSP, Utrecht (2000). 13. S.K. Duplock, J.G. Matisons, A.G. Swincer and R.F.O. Warren, J. Inorg. Organomet. Polym. 1, 361 (1991). 14. L.G. Britcher, D.C. Kehoe, J.G. Matisons and A.G. Swincer, Macromolecules, 28, 3110 (1995). 15. S. Hofmann, in Practical Surface Analysis, Second Edition, Vol 1 Auger and X-ray Photoelectron Spectroscopy, D. Briggs and M.P. Seah (Eds), John Wiley, New York (1990). 16. L.G. Britcher, J.G. Matisons and D.C. Kehoe, Unpublished results. 17. J.A. Brydson, Rubber Chemistry, Applied Science Publishers, London (1978). 18. E.P. Plueddemann, Silane Coupling Agents, 2nd ed, Plenum Press, New York (1991). 19. J. Blitz, R. Murthy and D. Leyden, J. Colloid Interface Sci. 126, 387 (1988). 20. A. Provatas and J.G. Matisons, Langmuir, 14, 1656 (1997). 21. D.W. Dwight, F.M. Fowkes and T. Huang, J. Adhesion Sci. Technol. 4, 619 (1990). 22. H. Ishida, in: Adhesion Aspects of Polymeric Coatings, K.L. Mittal (Ed.), pp 45-106, Plenum Press, New York (1983).



Surface modification of polyphenylene sulfide plastics to improve their adhesion to a dielectric adhesive YING WANG∗ and STANTON RAK Physical Technology Center, Automotive and Industrial Electronics Group, 4000 Commercial Avenue, Motorola, Inc, Northbrook, IL 60062

Abstract—The adhesion strength of a perfluorinated dielectric adhesive to polyphenylene sulfide (PPS) was investigated. The effect of different fillers in the PPS as a function of plasma treatment conditions was evaluated. The change in adhesion as a result of thermal baking was also addressed. The surface composition and surface energy were monitored and systematically quantified by X-ray Photoelectron Spectroscopy (XPS) and contact angle measurements, respectively. The correlations between the presence of certain functional groups, change in surface energy and polarity, and variation in adhesion properties indicate that the adhesion mechanism is mainly due to van der Waals forces. Enhanced wetting at the adhesive/substrate interface and a deeper interfacial diffusion zone are found to be necessary conditions to achieve the optimal adhesion. Keywords: Perfluorinated adhesive; surface modification; RF plasma; polyphenylene sulfide plastics.

1. INTRODUCTION

PPS is a widely used plastic in commercial and automotive industries. It has excellent mechanical strength and resistance to different media due to its structural stiffness provided by aromatic rings. This stiffness, however, sometimes prevents good adhesion of PPS to other materials. In order to improve the adhesion of PPS and other polymer materials, several surface modification techniques have been developed including chemical treatment, corona, arc discharge and plasma etching. Plasma surface treatment is an effective and clean way of improving the adhesion of many plastics. Several mechanisms have been proposed to interpret this improvement. They are: removal of organic contamination from the surfaces; material removal by ablation to increase the surface area or to remove a weak boundary layer; crosslinking or branching to strengthen the surface cohesively; and surface chemistry modification to improve chemical and physical interactions at the ∗

To whom all correspondence should be addressed. Phone: (847) 480-4718, Fax: (847) 480-3064, E-mail: [email protected]

Y. Wang and S. Rak

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bonding interface [1–3]. To determine the mechanisms behind the modification of the plastic, several surface sensitive techniques have been used including X-ray photoelectron spectroscopy (XPS) [4–6], infrared reflection absorption spectroscopy (IRRAS) [4], and contact angle measurements [2, 7]. It is generally accepted that the first step in the O2 plasma treatment consists of the formation of radicals on the surface layer by the elimination of hydrogen atoms, [8]. The radicals then react with oxygen and water molecules in air to produce functional groups such as alcohols, ketones and carboxylic acids, which are capable of interacting with an adhesive [7, 9]. It has also been reported that these functional groups normally have a short life-time because of the high mobility of plasma-induced, fragmented polymer chains on the surface. Therefore, to maintain good adhesion after plasma treatment, the effect of aging in different environmental conditions must be carefully monitored and controlled. It should be noted that depending on the nature of the bonding between the adhesive and substrate, the adhesion strength as well as the stability in different media, could be very different. In industrial applications, certain fillers including glass fibers and minerals are added into the pure polymer matrix to improve its thermal and mechanical properties. At the same time, the introduction of fillers may result in heterogeneity of surface morphology and chemistry and influence adhesion behavior and mechanisms compared to single phase plastics. The adhesive suppliers add adhesion promoters to promote adhesion with fillers such as SiO2. The mechanism for adhesion to PPS itself is less well understood and is of practical interest. This work describes the use of plasma treatment for improving the adhesion of a perfluorinated adhesive to PPS. The mechanism of this improvement was investigated using XPS and surface energy characterizations. PPS plastics with and without glass fillers and minerals were used. The effects of etching conditions and thermal annealing are also addressed. 2. EXPERIMENTAL DETAILS

2.1. Materials and sample preparation The chemical structure of PPS is shown below,

S n

polyphenylene sulfide (PPS) Three types of natural colored PPS samples from Ticona (Summit, NJ) were used to investigate the effects of fillers on adhesion. They were PPS without glass mineral additive (Fortronâ 0214 P1), PPS with glass fibers (Fortronâ 1140 L4),

Surface modification of polyphenylene sulfide plastics

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and PPS with glass fibers plus mineral fillers (Fortronâ 4665 B6). The weight percentages of SiO2 fibers and calcium-containing minerals in each sample, obtained by thermogravimetric analysis at 800 °C, are listed in Table 1. As shown in the table, the major components of the three materials are very different. These differences also imply that the materials are formulated under different conditions, and likely with other minor additives. As a result, material properties that are directly and indirectly related to adhesion properties, such as density, glass transition temperature, molecular weight, and molding temperature, can also be different. Due to this complexity, it is very difficult to unambiguously identify the differences in bonding characteristics between materials. Therefore, the emphasis in this work was placed on the mechanism of surface modification and its effects on adhesion within the same material group. The PPS samples in the form of tensile bars were first rinsed in methanol in an ultrasonic cleaner for 5 minutes to remove possible surface grease and low molecular weight contaminants. The thermal treatments after plasma exposure were conducted at 85 °C. Table 1. Compositions of PPS materials (wt%) Brand name

Plastic

Filler

Others additives

Fortronâ 0214 P1 (natural color)

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